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Nanoporous metals from thermal decomposition of transition metal dichalcogenides
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Nanoporous Metals from Thermal Decomposition of Transition Metal Dichalcogenides Swarnendu Chatterjee , Anton Anikin , Debjit Ghoshal , James L. Hart , Yawei Li , Saad Intikhab , D.A. Chareev , O.S. Volkova , A.S. Vasiliev , Mitra L. Taheri , Nikhil Koratkar , Goran Karapetrov , Joshua Snyder PII: DOI: Reference:
S1359-6454(19)30749-9 https://doi.org/10.1016/j.actamat.2019.11.018 AM 15644
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Acta Materialia
Received date: Revised date: Accepted date:
19 September 2019 6 November 2019 7 November 2019
Please cite this article as: Swarnendu Chatterjee , Anton Anikin , Debjit Ghoshal , James L. Hart , Yawei Li , Saad Intikhab , D.A. Chareev , O.S. Volkova , A.S. Vasiliev , Mitra L. Taheri , Nikhil Koratkar , Goran Karapetrov , Joshua Snyder , Nanoporous Metals from Thermal Decomposition of Transition Metal Dichalcogenides, Acta Materialia (2019), doi: https://doi.org/10.1016/j.actamat.2019.11.018
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Nanoporous Metals from Thermal Decomposition of Transition Metal Dichalcogenides
Swarnendu Chatterjeea, Anton Anikinb, Debjit Ghoshalc, James L. Hartd, Yawei Lia, Saad Intikhaba, D.A. Chareeve,f,g, O.S. Volkovah,i,j, A.S. Vasilieve,h,i, Mitra L. Taherid, Nikhil Koratkark, Goran Karapetrovb, and Joshua Snydera,* a
Department of Chemical and Biological Engineering, Drexel University, Philadelphia, PA
19104, USA b
c
Department of Physics, Drexel University, Philadelphia, PA 19104, USA
Department of Chemical and Biological Engineering, Rensselaer Polytechnic Institute, Troy,
NY 12180, USA d
Department of Materials Science and Engineering, Drexel University, 3141 Chestnut Street,
Philadelphia, Pennsylvania 19104, United States e
Ural Federal University, 620002 Ekaterinburg, Russia
f
Institute of Experimental Mineralogy, RAS, 142432 Chernogolovka, Russia
g
Institute of Geology and Petroleum Technologies, Kazan Federal University, Kazan, Russia
h
Low Temperature Physics and Superconductivity Department, M.V. Lomonosov Moscow State
University, Moscow 119991, Russia i
National University of Science and Technology MISiS, Moscow 119049, Russia
j
National Research South Ural State University, 454080 Chelyabinsk, Russia
k
Department of Mechanical, Aerospace and Nuclear Engineering, Rensselaer Polytechnic
Institute, Troy, NY 12180, USA *
Corresponding Author: Joshua Snyder, [email protected]
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Graphical abstract
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Abstract Nanoporous metals (np-M) have emerged as promising materials owing to their high surface area-to-volume ratio and electrical/thermal conductivity. There exists a group of processing methodologies by which np-M are formed through a top-down nanostructure evolution driven by the selective removal of a sacrificial component, all of which are a variation of dealloying. Nanoporosity evolution through current dealloying methodologies, however, is governed by strict requirements including sufficient separation in “reactivity” of the participating components and a homogeneous solid solution precursor alloy. This limits the viable alloy systems that may be used and the range of np-M’s that may be formed. Here, we report thermal decomposition of crystalline transition metal dichalcogenides (TMDs) as a new processing methodology for np-M formation, adding to the spectrum of dealloying protocols. We demonstrate application of this process to the formation of a broader class of np-M including W, Re, Mo, and Ta with feature sizes below 100 nm. The presented facile thermal treatment of TMDs offers a new methodology for the evolution of nanoporosity in a broad range of metals.
Keywords: Nanoporous metal, dealloying, nanoporosity, transition metal dichalcogenides
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1.
Introduction Nanoporous metals formed through alloy corrosion are characterized by a unique three
dimensional bicontinuous architecture, tunable pore size and composition, and high surface-tovolume ratio, giving them broad applicability for use as free-standing catalysts1–5, electronic devices6–11, sensor materials3,12, actuators13–15, etc. The nanoporosity evolution process, known as dealloying, is a top-down methodology involving the selective removal of a less noble or more reactive component from a binary or multicomponent alloy. The remaining metal component reorganizes, through a process that is well described by spinodal decomposition, to form a bicontinuous, three-dimensional nanostructure. To date there are four distinct methodologies for the formation of nanoporous metals: (1) electrochemical and chemical dealloying16,17,26,18–25, (2) liquid metal dealloying27–34, (3) evaporative dealloying35, and (4) reductive decomposition of ionic salts36,37. Electrochemical/chemical dealloying17,38 is the most recognized method to obtain nanoporosity and has resulted in the formation of nanoporous Pt4, Pd39, Au12, Ag40, Cu41, Ni42, Al26, Cu-Ni alloy43, etc. Layer-by-layer etching of the more reactive component initiates a pattern forming instability where the remaining more noble component diffuses against concentration gradients along the evolving surface to form an interconnected porous network22– 25
. Dealloying is a competition between dissolution and surface diffusion and the final
morphology/length scale is directly related to the ratio of the rates of these processes 22–25. Evolution of nanoporosity through electrochemical/chemical dealloying is predicated on two important precursor material criteria: (1) formation of a homogeneous, single-phase solid solution or intermetallic phase, and (2) a sufficient difference in the equilibrium redox potentials for the two metal components. These criteria severely limit the composition of the nanoporous
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metals that may be formed, mostly limiting it to noble metals that can survive the dealloying process
to
form
true
bicontinuous
porosity
at
the
nanoscale17,19,20,22–25.
Whereas
electrochemical/chemical dealloying is essentially driven by the difference in solubility of one metal ion over another in an aqueous/nonaqueous electrolyte, liquid metal dealloying is driven by the difference in solubility of the constituent metals in the alloy in a liquid metal melt27–34. With this technique, nanoporous architectures composed of metals and metalloids that were not previously attainable with electrochemical/chemical dealloying could be produced, as seen for Ta29 and Si34. The elevated temperatures required for the liquid metal melt “electrolyte”, however, typically limit the accessible pore sizes to greater than 100 nm27–34. Additionally, a final chemical/electrochemical step is still required to remove the solidified liquid metal “electrolyte” and expose the nanoporous structure29,44. More recently, vapor phase dealloying35 (thermal evaporation of high vapor pressure transition metals under high vacuum conditions) has been identified as a methodology for nanoporous metal synthesis. Here the driving force for dealloying is the difference in the rate of evaporation, established by their vapor pressures, between the two constituent metals making up the alloy. Nanoporous morphology and pore size can be controlled through the manipulation of the dealloying temperature and pressure. Vapor phase dealloying has been used to synthesize nanoporous Co, Si, Ti, and Ni35. This technique, however, is still limited by constraining intrinsic material properties. Analogous to electrochemical/chemical dealloying, a sufficient difference in vapor pressure of the alloy components is required to drive removal of one element and curvature driven diffusive rearrangement of the other. Additionally, for true bicontinuous nanoporosity, the alloying components must form a single phase, homogeneously distributed alloy. Both of these constraints significantly limit the breadth of elements that may be used for this process. So far,
5
Zn has been the only viable sacrificial element due to its high vapor pressure. Finally, reductive decomposition of ionic metal salts has demonstrated success for the formation of nanoporous Ag36, Cu37, Ni45 and Sn37. However, pore size and morphology control are limited to larger pore diameters, >100 nm37,45. The limitations imparted by the strict precursor material property requirements constrain the applicability of the aforementioned dealloying protocols, particularly when trying to reduce nanostructure length scales below 100 nm for the more reactive and higher melting metals. Herein, we present thermal decomposition of transition metal dichalcogenides (TMDs) (sulfides, selenides, and mixed sulfo-selenides) as a facile procedure for the formation of nanoporous metals that are not easily synthesized through the previously listed dealloying techniques. Evaporative/reactive removal of the chalcogen atoms from TMDs, which we designate as thermal decomposition, results in the evolution of classical nanoporosity with pore sizes below 100 nm. We demonstrate the synthesis of nanoporous W, Re, Ta, Mo, Ti, Co, and Ni through the thermal decomposition of crystalline TMD precursor materials. Analysis of the effect of reactive gas on evolved morphology indicates that chalcogen loss occurs through a combination of evaporation and reactive conversion to H2X (X = S, Se). Additionally, the initial crystal structure, layered vs. pyrite-type, of the TMD is shown to have a significant effect on both the final nanoporous morphology and residual composition. The high vapor pressure and reactivity of chalcogen atoms along with the broad library of accessible TMD compositions, make this technique ideal for synthesis of a larger class of nanoporous metals that are not attainable through previously developed methodologies.
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2.
Experimental
2.1
Precursor Material Synthesis
Pyrite-type Co and Ni disulfide, diselenide, and sulfoselenide TMDs were synthesized through growth in eutectic mixtures of molten alkali salts in a stationary temperature gradient. For more details see refs.46–48. Layered, van der Waals TMDs with transition metals Ti, Ta, and W were prepared through traditional chemical vapor deposition processes49–51. Briefly, powders of pure metal and chalcogen along with iodine, which serves as a transport gas, are sealed in long quartz ampules at 10 – 20 mTorr. The sealed quartz tubes are placed in a 3-zone furnace at a desired temperature/temperature gradient for approximately two weeks. Upon the completion of crystal growth, pyrite-type crystals were removed from the solidified melt by dissolution in a distilled water, alcohol, and acetone mixture using an ultrasonic cleaner. The product was then dried in a muffle furnace at 70 oC. All the precursors used in the above synthesis were purchased as powders from Alfa Aesar with purity >99.99%. For Mo and Re, commercially available MoSe2 and ReSe2 were used (hqgraphene, purity > 99.995%). 2.2
Thermal Decomposition of TMDs and np-M formation
Thermal decomposition of TMDs including WSe2, WS2, ReSe2, TaS2, MoSe2, TiSe2, NiSe2, CoS2, CoSe2, and CoSxSe2-x was performed at various temperatures under flowing H2(5%)/Ar or Ar in a split-hinge tube furnace, schematically represented in Figure 1(e). Prior to thermal decomposition, the tube is purged with H2(5%)/Ar or Ar (10 psi pressure) at room temperature for 15 minutes to ensure complete removal of oxygen. Subsequently, the tube is quickly placed inside the tube furnace already heated to the desired temperature and the sample is thermally decomposed for the desired amount of time followed by quenching in ambient air.
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2.3
Material Characterization
The atomic compositions of the TMDs were confirmed in a digital scanning electron microscope TESCAN Vega II XMU with energy dispersive microanalysis system INCA Energy 450/XT (20 kV). The chemical compositions of the TMDs were shown to match the ratio of initially charged precursors. No significant presence of impurity elements belonging to the precursor salts was found. Raman measurements were taken at room temperature using a commercial WITEC 300 R confocal imaging setup. The spot size of ~ 2 µm-diameter was attained using a 50X objective. Raman signals were collected with the same objective and measured with a single-grating spectrometer and a cooled CCD camera. The nanoporous morphologies were analyzed with scanning electron microscope (SEM, Zeiss Supra 50VP) operated at 4 kV and transmission electron microscope (TEM, JEOL JEM-2100) operated at 200 kV. For in-situ TEM, a Gatan 628 single-tilt heating stage was used with a heating rate of 50 oC/minute. X-ray diffraction (XRD) spectra were analyzed on a Rigaku SmartLab diffractometer with a Cu Kα X-ray source operated at 40 kV and 30 mA. 3.
Results and Discussion
3.1
Thermal decomposition of TMDs
Representative samples of the nanoporous morphology that evolves through the thermal decomposition of crystalline TMDs is shown in Figure 1(a)-(c). The nanoporosity is found to be more than a superficial pore structure on the surface of the TMD crystals. A visual inspection of a focused ion beam (FIB) cross-section, Figure 1(d), shows homogeneous, bicontinuous nanoporosity at depths that are orders of magnitude larger than the characteristic length scale of the nanoporous metal. The FIB cross-section in Figure 1(d) shows a depth of porosity of ~5 µm; 8
however, there is nothing to suggest that this methodology cannot produce nanoporosity further into the bulk of the material. Traditional dealloying is driven by the solubilization of a more reactive component into an external media and the subsequent rearrangement of the exposed, lower coordinated, less reactive component. Here, elevated temperature and reactive gas is used to drive the loss of the chalcogen elements. Corollaries can be drawn to the recent report on vapor-phase dealloying in which high vapor pressure metals, such as Zn, are evaporatively removed from a homogeneous alloy35. The breadth of nanoporous metals that can be made through vapor-phase dealloying, however, is limited. TMD precursors, in contrast, present a wide breadth of available transition metals52, both single metal and mixed metals, in an even and homogeneous compositional distribution within well-defined crystalline structures. Chalcogen elements (S, Se, Te) exhibit high vapor pressures and are susceptible to reactive etching. These qualities make it easier to find processing conditions that yield nanoporosity evolution without adversely impacting the stability of the transition metal, especially for sub 100 nm porous refractory metals. We demonstrate the broad applicability of thermal decomposition of TMDs for nanoporous metal formation through the observation of nanoporosity evolution from both layered and pyrite-type precursor MX2 crystal structures with X = S, Se: np-W (Figures 1, 5, 2, S4, S5, S6, S9), np-Re (Figures 1, S4, and S9), np-Mo (Figures 1 and S9), np-Co (Figure 6, S1, and S4), np-Ni (Figure S1), np-Ti (Figure S1), and np-Ta (Figure S1). This is particularly remarkable as formation of nanoporous metals composed of high melting, refractory elements through conventional dealloying methodologies is likely to either be stymied by oxidative passivation at the etch front, or limited to length scales considerably larger than 100 nm27–34. Thermally driven porosity formation in TMDs propagates through two mechanisms that operate in parallel: (1) reactive formation of gaseous H2X due to the thermally induced reaction
9
M (s) H 2 X ( g ) ), where X = S/Se and M is the metal) and with hydrogen ( MX (s) H 2 ( g )
(2) evaporation of chalcogen atoms from the TMD surface. Support for this mechanism is shown in Figure 2 comparing the morphology of thermally decomposed WSe2 at 950 °C in pure Ar (evaporative chalcogen removal only) to that in the presence of H2 (H2(5%)/Ar). The sample decomposed in H2/Ar is found to have more homogeneous porosity, while that in Ar has a discontinuous, bimodal porosity. Additionally, the final morphology length scale is smaller in H2/Ar than in Ar, 25 nm and 48 nm respectively, for the same decomposition conditions, 950 oC and 1 hr. This result indicates the reaction with H2 aids the fast removal of chalcogen atoms and consequently smaller initial porosity length scale. This behavior is also seen in classical electrochemical dealloying where higher rates of corrosion shift the balance between dissolution and surface diffusion, nucleating smaller length scales at early times22,53. Each of the two processes (evaporative and reactive) initiates the loss of chalcogen atoms through covalent bond severance at the gas-solid interface. This process is analogous to the oxidation and formation of a soluble cation in the electrochemical/chemical dealloying mechanism. As with traditional dealloying methodologies, the initiating event is the formation of a vacancy in the TMD lattice through the removal of a chalcogen atom. The exposure of low coordinated transition metal atoms on the surface induces a surface smoothening driven diffusion, mimicking spinodal decomposition of adatoms and vacancies, and formation of a porous architecture. The resulting morphology, as shown in Figure 1, resembles that of the bicontinuous nanoporosity observed in traditional dealloying systems16–25. 3.2
Effect of TMD precursor crystal structure
The precursor TMD crystal structure is found to have a significant effect on both the mechanism of porosity evolution and the final morphology/composition attained after processing. np-Ni and 10
np-Co are obtained through the thermal decomposition of pyrite-type TMDs54, Figure S1. These TMDs are defined by a long range crystal structure, covalent bonding throughout, where the metal atom is surrounded by six neighboring chalcogen atoms in a distorted octahedral orientation46,54,55. This more closely resembles the fcc structure of metal alloys in a traditional dealloying system. In contrast, layered TMDs possess a layered crystal structure where each MX2 layer (M = W, Re, Mo, Ti, Ta; X = Se, Se) is only connected to the neighboring layers through van der Waals attraction52,56. To understand the effect of pyrite-type vs. layered crystal structures, we compare the formation of np-Co and np-W. XRD, Figure 3(a), for the np-Co and np-W pre- and post-porosity evolution indicate the appearance of metallic phases of Co(002), CoSe(101), and CoSe(110) in np-Co and W(110) phase in np-W along with some residual TMD phases in the nanoporous material. The W(110) phase in np-W is also confirmed from the measured lattice parameter (0.2223 nm), Figure S2. Figure 3(b) shows the residual chalcogen contents in np-Co and np-W as a function of thermal decomposition time. For np-Co, the chalcogen content shows a steady decrease with time, approaching 25 at.% X (X = S, Se) after 9 hours. For np-W, the chalcogen content drops significantly in the first hour of decomposition. After 9 hours, np-W is found to be essentially pure W, with nearly all of the chalcogen removed. The time dependent loss of chalcogen content is confirmed through Raman spectroscopy (Figure S3). The change in A1g and E2g Raman modes57 of WSe2 are measured at different decomposition times. Both the out of plane A1g mode and in-plane E2g modes of WSe2 are visible up to 1 hour but disappear after 3 hours. Observation of the morphology and composition as it evolves with time and temperature for the pyrite-type, Co-based TMDs to np-Co shows a morphology and residual composition very similar to that of traditional dealloying, Figures 3, 6, and S1(b). The island formation and undercutting that define the early stages of dealloying, Figure 4(b), and
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etch front propagation result in the formation of a core-shell structure, retaining a high fraction of the more reactive, less noble component within the interior of the ligaments4,5,19,20,22. In contrast to metal alloys and pyrite-type TMDs, only weak van der Waals forces hold MX2 layers together in layered TMDs (where M = W, Mo, Ti, Re etc.). This has two effects: (1) following chalcogen removal, there is no direct route, at the single atom scale, for the formation of a metal-metal bond between the transition metals in neighboring layers; (2) there is the potential for reactive gas intercalation between the MX2 layers. Regarding the latter, layered TMDs have been shown to be able to absorb/intercalate hydrogen, to the point where they have been evaluated as hydrogen storage materials58–61. This intercalation of reactive hydrogen facilitates chalcogen removal from both the exterior of the crystalline structure as well as the interior. The consequent result is the fast and complete reactive removal of the chalcogen element throughout the crystalline material, not simply limited to an etch front moving into the bulk of the crystal from a surface. Raman and EDS elemental analysis shows that thermal decomposition of the layered TMDs results in near complete chalcogen element removal, Figures 3 and S3. The interlayer chalcogen loss, to the point of full removal, along with the absence of any obvious mechanism for the formation of metal-metal bonds between the single atomic plane of metals in neighboring TMDs sheets suggests that the traditional dealloying mechanism is not operative here. Surface morphology at early times, 5 minutes of thermal decomposition, Figure 4, for the layered and pyrite-type TMDs show a distinct difference. For pyrite-type TMDs, classical mounding is observed and the initiation of the ligamentous structure is apparent. For the layered TMDs, however, the surface is decorated with small particles. With this observation, taken with the compositional data and the fact that H2 can intercalate between the layers, we can put forward a proposed mechanism for the formation of interconnected
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nanoporosity through the thermal decomposition of layered TMDs. As a consequence of the absence of metal-metal bonding between the MX2 layers and the intercalation of H2, the removal of chalcogen from the MX2 layers results in the coalescence of metal atoms into nanoscale clusters during thermal decomposition of the individual TMD layers. Occurring simultaneously internally and externally, the interconnected, bicontinuous nanoporosity then evolves through the aggregation and curvature driven smoothening of the aggregate structure as particles formed on adjoining layers coalesce. The observation of a small fraction of intra-ligament grain boundaries, Figure S4(a) and (b), for np-M formed from layered TMDs which are absent for np-M formed from the pyrite-type TMDs, Figure S4(c) and (d), supports this hypothesis. Figure 5 is a series of bright field TEM images taken during in-situ heating of WSe2 at 950 oC in an inert environment (10-5 Pa). This series of images shows the early stages of feature formation, initiating at the edges and propagating inward. Complete bicontinuous nanoporosity is attained after ~7 minutes. In-situ HRTEM video (Supporting Information), from which the still images in Figure 5 are obtained, shows an apparent uniform rate of etch front propagation and ligament formation progressing from the surface towards the bulk. This indicates that porosity evolution is likely interface reaction limited, which consists of chalcogen removal followed by transition metal diffusion and aggregation, and vapor phase chalcogen removal through both the pores and between the TMD layers is sufficiently fast. Further evidence for an interface limited process is presented in Figure 1(d) where there is an absence of a gradient of ligament size for porosity depths on the order of microns. The end result of the process is a sub-100 nm, uniform bicontinuous nanoporous metal with porosity depth on the order of microns (Figure 1(d)) for refractory metals (Re, Mo, W, Ta, Ti) which to this point has been very difficult to achieve otherwise. For these refractory metals, the degree of post-porosity evolution coarsening, even at
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temperatures as high as 950 °C, is low due to the slow rate of surface diffusion. Surface diffusion rates of transition metals roughly scale inversely with the metal’s melting point62 . The np-W pore size remains below 100 nm even after 24 hr at 950 oC, Figure S5. 3.3
Rate controlling process: “Dissolution” vs. Diffusion
For nanoporosity formation through the removal of a sacrificial species, it is easy to think of the “corrosion” event as the defining process for both the rate of porosity evolution and the final morphology. Rate of removal of the more reactive chalcogen component of the TMDs is related to the composition of the TMD through the bond strength between the transition metal and chalcogen as well as the temperature dependence of the evaporation rate, related to the vapor pressure of the chalcogen and/or its reactivity with H2. In Figure 3(c) and (d)63, the experimentally observed minimum temperature required to initiate chalcogen loss and propagate nanoporosity is plotted versus both the melting point of the transition metal in the TMD as well as the formation energy of MS2 TMDs. A strong correlation is observed between the temperature required to initiate nanoporosity formation and the melting point of the transition metal within the TMD, Figure 3(c). In contrast, no significant correlation is observed with the formation energy of MX2. Experimentally we find that the identity of the chalcogen (X = S, Se) in layered TMDs, Figures S6, has no effect on the final morphology. Conversely, a modest effect of chalcogen identity is observed for the pyrite-type TMDs, in terms of average pore size, as a function of thermal decomposition temperature, Figure S7, and at a range of CoSxSe2-x compositions spanning from CoS2 to CoSe2, Figure 6(a)-(e). The M-Se bond is known to be stronger than the M-S bond46,64 and the vapor pressure of S is higher than that of Se, Figure 6(f). Therefore, MS2 should have a higher rate of thermal decomposition in comparison to MSe2 TMDs. The absence of any visible effect of chalcogen identity for the layered TMDs stems from 14
the low rate of surface diffusion for the high melting refractory metals from which they are composed. This low rate of surface reorganization dominates over any effects associated with chalcogen removal. The pyrite-type TMDs tested here are composed of transition metals with lower melting points, relative to the refractory metal containing layered TMDs, which corresponds to higher rates of surface diffusion at a given temperature. We argue that the observed effect of chalcogen identity for the pyrite-type TMDs, however modest, is due in part to the shift in the balance of the competition between “dissolution” and diffusion as a consequence of the higher rates of surface diffusion for the lower melting metals. The strong correlation between metal melting point and the minimum temperature required for nanoporosity formation in Figure 3(c) and the observed range in sensitivity to the chalcogen identity, from negligible for layered TMDs of higher melting transition metals to moderate for pyrite-type TMDs of lower melting transition metals, indicates that porosity formation and propagation is dominated by the surface diffusion of the remaining transition metal exposed through the removal of the chalcogen element. Activation energies for the thermally driven porosity evolution in CoS2 and CoSe2 are determined by plotting average ligament diameters versus temperature at early decomposition times (30 min.), Figure S8. The porosity evolution activation energies for CoS2 and CoSe2 are 99.96 kJ/mol (1.036 eV) and 104.02 kJ/mol (1.078 eV) respectively. Thermal evaporation of Zn from CoZn alloys was measured to have an activation barrier of 2.28 eV at low vacuum and 0.25 eV at high vacuum35. The difference in activation barrier was attributed to the shift in limiting processes. At low vacuum, the evaporation of Zn, speculated to be controlled by the bulk diffusion of Zn, is limiting. At high vacuum, the surface diffusion of Co is limiting35. The values obtained here lie within this range and suggest that porosity evolution is controlled by surface
15
diffusion of the transition metal. The absolute value of the activation barriers determined for CoS2 and CoSe2 thermal decomposition at atmospheric pressure, being closer to those determined at high vacuum conditions for vapor phase dealloying, is due to the reactive etching of the chalcogen elements in addition to simple evaporation. While the value of the activation barrier is suggestive of a surface diffusion limited process, there is a slight compositional dependence to the value, in line with data in Figure 6 and S7. This could be related to the slight difference in the rate of M-X bond breaking as the M-Se bond is stronger than the M-S bond46,64 and the vapor pressure of S is higher than Se, Figure 6(f). Overall, however, these results indicate a surface diffusion-controlled process, in agreement with other dealloying systems and methodologies, as well as the other data presented in this study. 4.
Conclusions In summary, we present a new, broadly applicable protocol for the synthesis of
morphologically controlled np-M with length scales below 100 nm, based on thermal decomposition of TMDs. With this new methodology, we have successfully synthesized a broad range of np-M where M = W, Re, Ta, Mo, Ti, Co, and Ni. Analysis of the impact of reactive gas, precursor crystal structure, and TMD composition gives insight into the mechanisms by which nanoporosity is evolved. Thermal decomposition offers a novel and universal pathway to synthesize a broad range of np-M from a diverse library of available TMDs.
Declaration of Competing Interest None.
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23
Figures
Figure 1: Scanning electron micrographs (SEM) of (a) np-W formed through the thermal decomposition of layered WSe2 at 950 °C under flowing H2(5%)/Ar (scale bar: 100 nm), (b) npMo: MoS2 precursor, layered, 800 °C, H2(5%)/Ar, 0.5 hrs. (scale bar: 100 nm), and (c) np-Re: ReS2 precursor, layered, 950 °C, H2(5%)/Ar, 0.5 hrs. (scale bar: 100 nm). (d) Focused ion beam (FIB) cross section of np-W (scale bar: 1 µm). (e) Schematic representation of the thermal decomposition process for the evolution of nanoporosity in TMDs.
24
Figure 2: Demonstration of the effect of gaseous atmosphere (Ar vs. H2/Ar) on thermal decomposition of TMDs. (a) and (b) np-W from WSe2 after 1 hr. in Ar at 950 °C (scale bar: (a) 200 nm, (b) 2 µm). (c) and (d) np-W from WSe2 after 1 hr. in H2(5%)/Ar at 950 °C (scale bar: (a) 200 nm, (b) 1 µm). Images (a) and (c) are at the same magnification.
25
Figure 3: (a) XRD spectra of CoSe2 (blue), np-Co (purple), WSe2 (black), and np-W (red). (b) Chalcogen content measured using SEM energy dispersive spectroscopy (EDS) as a function of thermal decomposition time for layered WSe2 (black) and pyrite-type CoSe2 (blue) crystalline precursor materials. (c) Minimum temperature required to initiate porosity formation through chalcogen removal and remaining metal surface diffusion in MX2 TMDs (X = S, Se; M = Ni, Co, Ti, Mo, Ta, Re, W) plotted versus the metal point of the pure transition metal. (d) Minimum temperature required to initiate porosity formation through chalcogen removal and remaining metal surface diffusion in MX2 TMDs plotted versus the formation energy of MS2. Red lines in (c) and (d) are linear fits to the data points.
26
Figure 4: (a) SEM of the early stages of nanoporous morphology formation in WSe2 through thermal decomposition (5 min., 950 °C, H2(5%)/Ar), scale bar: 100 nm. (b) SEM of the early stages of nanoporous morphology formation in CoSe2 through thermal decomposition (5 min., 500 °C, H2(5%)/Ar), scale bar: 200 nm.
27
Figure 5: In-situ heated stage TEM observing the time dependent morphology during thermal decomposition of WSe2 at 950 °C. (a) 0 minutes, (b) 4 minutes, (c) 5 minutes, (d) 7 minutes. Scale bars are 20 nm.
28
Figure 6: SEMs of np-Co formed through the thermal decomposition (500 °C, 20 min, H2(5%)/Ar) of CoS(x)Se(2-x) with (a) x = 2, (b) x = 1.4, (c) x = 0.6, (d) x = 0. Scale bars are 1 µm. (e) Average pore size as a function of composition and thermal decomposition time; 20 min (black) and 30 min (red). (f) Vapor pressure as a function of temperature for S (blue) and Se (green).
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Nanoporous Metals from Thermal Decomposition of Transition Metal Dichalcogenides Swarnendu Chatterjee , Anton Anikin , Debjit Ghoshal , James L. Hart , Yawei Li , Saad Intikhab , D.A. Chareev , O.S. Volkova , A.S. Vasiliev , Mitra L. Taheri , Nikhil Koratkar , Goran Karapetrov , Joshua Snyder PII: DOI: Reference:
S1359-6454(19)30749-9 https://doi.org/10.1016/j.actamat.2019.11.018 AM 15644
To appear in:
Acta Materialia
Received date: Revised date: Accepted date:
19 September 2019 6 November 2019 7 November 2019
Please cite this article as: Swarnendu Chatterjee , Anton Anikin , Debjit Ghoshal , James L. Hart , Yawei Li , Saad Intikhab , D.A. Chareev , O.S. Volkova , A.S. Vasiliev , Mitra L. Taheri , Nikhil Koratkar , Goran Karapetrov , Joshua Snyder , Nanoporous Metals from Thermal Decomposition of Transition Metal Dichalcogenides, Acta Materialia (2019), doi: https://doi.org/10.1016/j.actamat.2019.11.018
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Nanoporous Metals from Thermal Decomposition of Transition Metal Dichalcogenides
Swarnendu Chatterjeea, Anton Anikinb, Debjit Ghoshalc, James L. Hartd, Yawei Lia, Saad Intikhaba, D.A. Chareeve,f,g, O.S. Volkovah,i,j, A.S. Vasilieve,h,i, Mitra L. Taherid, Nikhil Koratkark, Goran Karapetrovb, and Joshua Snydera,* a
Department of Chemical and Biological Engineering, Drexel University, Philadelphia, PA
19104, USA b
c
Department of Physics, Drexel University, Philadelphia, PA 19104, USA
Department of Chemical and Biological Engineering, Rensselaer Polytechnic Institute, Troy,
NY 12180, USA d
Department of Materials Science and Engineering, Drexel University, 3141 Chestnut Street,
Philadelphia, Pennsylvania 19104, United States e
Ural Federal University, 620002 Ekaterinburg, Russia
f
Institute of Experimental Mineralogy, RAS, 142432 Chernogolovka, Russia
g
Institute of Geology and Petroleum Technologies, Kazan Federal University, Kazan, Russia
h
Low Temperature Physics and Superconductivity Department, M.V. Lomonosov Moscow State
University, Moscow 119991, Russia i
National University of Science and Technology MISiS, Moscow 119049, Russia
j
National Research South Ural State University, 454080 Chelyabinsk, Russia
k
Department of Mechanical, Aerospace and Nuclear Engineering, Rensselaer Polytechnic
Institute, Troy, NY 12180, USA *
Corresponding Author: Joshua Snyder, [email protected]
1
Graphical abstract
2
Abstract Nanoporous metals (np-M) have emerged as promising materials owing to their high surface area-to-volume ratio and electrical/thermal conductivity. There exists a group of processing methodologies by which np-M are formed through a top-down nanostructure evolution driven by the selective removal of a sacrificial component, all of which are a variation of dealloying. Nanoporosity evolution through current dealloying methodologies, however, is governed by strict requirements including sufficient separation in “reactivity” of the participating components and a homogeneous solid solution precursor alloy. This limits the viable alloy systems that may be used and the range of np-M’s that may be formed. Here, we report thermal decomposition of crystalline transition metal dichalcogenides (TMDs) as a new processing methodology for np-M formation, adding to the spectrum of dealloying protocols. We demonstrate application of this process to the formation of a broader class of np-M including W, Re, Mo, and Ta with feature sizes below 100 nm. The presented facile thermal treatment of TMDs offers a new methodology for the evolution of nanoporosity in a broad range of metals.
Keywords: Nanoporous metal, dealloying, nanoporosity, transition metal dichalcogenides
3
1.
Introduction Nanoporous metals formed through alloy corrosion are characterized by a unique three
dimensional bicontinuous architecture, tunable pore size and composition, and high surface-tovolume ratio, giving them broad applicability for use as free-standing catalysts1–5, electronic devices6–11, sensor materials3,12, actuators13–15, etc. The nanoporosity evolution process, known as dealloying, is a top-down methodology involving the selective removal of a less noble or more reactive component from a binary or multicomponent alloy. The remaining metal component reorganizes, through a process that is well described by spinodal decomposition, to form a bicontinuous, three-dimensional nanostructure. To date there are four distinct methodologies for the formation of nanoporous metals: (1) electrochemical and chemical dealloying16,17,26,18–25, (2) liquid metal dealloying27–34, (3) evaporative dealloying35, and (4) reductive decomposition of ionic salts36,37. Electrochemical/chemical dealloying17,38 is the most recognized method to obtain nanoporosity and has resulted in the formation of nanoporous Pt4, Pd39, Au12, Ag40, Cu41, Ni42, Al26, Cu-Ni alloy43, etc. Layer-by-layer etching of the more reactive component initiates a pattern forming instability where the remaining more noble component diffuses against concentration gradients along the evolving surface to form an interconnected porous network22– 25
. Dealloying is a competition between dissolution and surface diffusion and the final
morphology/length scale is directly related to the ratio of the rates of these processes 22–25. Evolution of nanoporosity through electrochemical/chemical dealloying is predicated on two important precursor material criteria: (1) formation of a homogeneous, single-phase solid solution or intermetallic phase, and (2) a sufficient difference in the equilibrium redox potentials for the two metal components. These criteria severely limit the composition of the nanoporous
4
metals that may be formed, mostly limiting it to noble metals that can survive the dealloying process
to
form
true
bicontinuous
porosity
at
the
nanoscale17,19,20,22–25.
Whereas
electrochemical/chemical dealloying is essentially driven by the difference in solubility of one metal ion over another in an aqueous/nonaqueous electrolyte, liquid metal dealloying is driven by the difference in solubility of the constituent metals in the alloy in a liquid metal melt27–34. With this technique, nanoporous architectures composed of metals and metalloids that were not previously attainable with electrochemical/chemical dealloying could be produced, as seen for Ta29 and Si34. The elevated temperatures required for the liquid metal melt “electrolyte”, however, typically limit the accessible pore sizes to greater than 100 nm27–34. Additionally, a final chemical/electrochemical step is still required to remove the solidified liquid metal “electrolyte” and expose the nanoporous structure29,44. More recently, vapor phase dealloying35 (thermal evaporation of high vapor pressure transition metals under high vacuum conditions) has been identified as a methodology for nanoporous metal synthesis. Here the driving force for dealloying is the difference in the rate of evaporation, established by their vapor pressures, between the two constituent metals making up the alloy. Nanoporous morphology and pore size can be controlled through the manipulation of the dealloying temperature and pressure. Vapor phase dealloying has been used to synthesize nanoporous Co, Si, Ti, and Ni35. This technique, however, is still limited by constraining intrinsic material properties. Analogous to electrochemical/chemical dealloying, a sufficient difference in vapor pressure of the alloy components is required to drive removal of one element and curvature driven diffusive rearrangement of the other. Additionally, for true bicontinuous nanoporosity, the alloying components must form a single phase, homogeneously distributed alloy. Both of these constraints significantly limit the breadth of elements that may be used for this process. So far,
5
Zn has been the only viable sacrificial element due to its high vapor pressure. Finally, reductive decomposition of ionic metal salts has demonstrated success for the formation of nanoporous Ag36, Cu37, Ni45 and Sn37. However, pore size and morphology control are limited to larger pore diameters, >100 nm37,45. The limitations imparted by the strict precursor material property requirements constrain the applicability of the aforementioned dealloying protocols, particularly when trying to reduce nanostructure length scales below 100 nm for the more reactive and higher melting metals. Herein, we present thermal decomposition of transition metal dichalcogenides (TMDs) (sulfides, selenides, and mixed sulfo-selenides) as a facile procedure for the formation of nanoporous metals that are not easily synthesized through the previously listed dealloying techniques. Evaporative/reactive removal of the chalcogen atoms from TMDs, which we designate as thermal decomposition, results in the evolution of classical nanoporosity with pore sizes below 100 nm. We demonstrate the synthesis of nanoporous W, Re, Ta, Mo, Ti, Co, and Ni through the thermal decomposition of crystalline TMD precursor materials. Analysis of the effect of reactive gas on evolved morphology indicates that chalcogen loss occurs through a combination of evaporation and reactive conversion to H2X (X = S, Se). Additionally, the initial crystal structure, layered vs. pyrite-type, of the TMD is shown to have a significant effect on both the final nanoporous morphology and residual composition. The high vapor pressure and reactivity of chalcogen atoms along with the broad library of accessible TMD compositions, make this technique ideal for synthesis of a larger class of nanoporous metals that are not attainable through previously developed methodologies.
6
2.
Experimental
2.1
Precursor Material Synthesis
Pyrite-type Co and Ni disulfide, diselenide, and sulfoselenide TMDs were synthesized through growth in eutectic mixtures of molten alkali salts in a stationary temperature gradient. For more details see refs.46–48. Layered, van der Waals TMDs with transition metals Ti, Ta, and W were prepared through traditional chemical vapor deposition processes49–51. Briefly, powders of pure metal and chalcogen along with iodine, which serves as a transport gas, are sealed in long quartz ampules at 10 – 20 mTorr. The sealed quartz tubes are placed in a 3-zone furnace at a desired temperature/temperature gradient for approximately two weeks. Upon the completion of crystal growth, pyrite-type crystals were removed from the solidified melt by dissolution in a distilled water, alcohol, and acetone mixture using an ultrasonic cleaner. The product was then dried in a muffle furnace at 70 oC. All the precursors used in the above synthesis were purchased as powders from Alfa Aesar with purity >99.99%. For Mo and Re, commercially available MoSe2 and ReSe2 were used (hqgraphene, purity > 99.995%). 2.2
Thermal Decomposition of TMDs and np-M formation
Thermal decomposition of TMDs including WSe2, WS2, ReSe2, TaS2, MoSe2, TiSe2, NiSe2, CoS2, CoSe2, and CoSxSe2-x was performed at various temperatures under flowing H2(5%)/Ar or Ar in a split-hinge tube furnace, schematically represented in Figure 1(e). Prior to thermal decomposition, the tube is purged with H2(5%)/Ar or Ar (10 psi pressure) at room temperature for 15 minutes to ensure complete removal of oxygen. Subsequently, the tube is quickly placed inside the tube furnace already heated to the desired temperature and the sample is thermally decomposed for the desired amount of time followed by quenching in ambient air.
7
2.3
Material Characterization
The atomic compositions of the TMDs were confirmed in a digital scanning electron microscope TESCAN Vega II XMU with energy dispersive microanalysis system INCA Energy 450/XT (20 kV). The chemical compositions of the TMDs were shown to match the ratio of initially charged precursors. No significant presence of impurity elements belonging to the precursor salts was found. Raman measurements were taken at room temperature using a commercial WITEC 300 R confocal imaging setup. The spot size of ~ 2 µm-diameter was attained using a 50X objective. Raman signals were collected with the same objective and measured with a single-grating spectrometer and a cooled CCD camera. The nanoporous morphologies were analyzed with scanning electron microscope (SEM, Zeiss Supra 50VP) operated at 4 kV and transmission electron microscope (TEM, JEOL JEM-2100) operated at 200 kV. For in-situ TEM, a Gatan 628 single-tilt heating stage was used with a heating rate of 50 oC/minute. X-ray diffraction (XRD) spectra were analyzed on a Rigaku SmartLab diffractometer with a Cu Kα X-ray source operated at 40 kV and 30 mA. 3.
Results and Discussion
3.1
Thermal decomposition of TMDs
Representative samples of the nanoporous morphology that evolves through the thermal decomposition of crystalline TMDs is shown in Figure 1(a)-(c). The nanoporosity is found to be more than a superficial pore structure on the surface of the TMD crystals. A visual inspection of a focused ion beam (FIB) cross-section, Figure 1(d), shows homogeneous, bicontinuous nanoporosity at depths that are orders of magnitude larger than the characteristic length scale of the nanoporous metal. The FIB cross-section in Figure 1(d) shows a depth of porosity of ~5 µm; 8
however, there is nothing to suggest that this methodology cannot produce nanoporosity further into the bulk of the material. Traditional dealloying is driven by the solubilization of a more reactive component into an external media and the subsequent rearrangement of the exposed, lower coordinated, less reactive component. Here, elevated temperature and reactive gas is used to drive the loss of the chalcogen elements. Corollaries can be drawn to the recent report on vapor-phase dealloying in which high vapor pressure metals, such as Zn, are evaporatively removed from a homogeneous alloy35. The breadth of nanoporous metals that can be made through vapor-phase dealloying, however, is limited. TMD precursors, in contrast, present a wide breadth of available transition metals52, both single metal and mixed metals, in an even and homogeneous compositional distribution within well-defined crystalline structures. Chalcogen elements (S, Se, Te) exhibit high vapor pressures and are susceptible to reactive etching. These qualities make it easier to find processing conditions that yield nanoporosity evolution without adversely impacting the stability of the transition metal, especially for sub 100 nm porous refractory metals. We demonstrate the broad applicability of thermal decomposition of TMDs for nanoporous metal formation through the observation of nanoporosity evolution from both layered and pyrite-type precursor MX2 crystal structures with X = S, Se: np-W (Figures 1, 5, 2, S4, S5, S6, S9), np-Re (Figures 1, S4, and S9), np-Mo (Figures 1 and S9), np-Co (Figure 6, S1, and S4), np-Ni (Figure S1), np-Ti (Figure S1), and np-Ta (Figure S1). This is particularly remarkable as formation of nanoporous metals composed of high melting, refractory elements through conventional dealloying methodologies is likely to either be stymied by oxidative passivation at the etch front, or limited to length scales considerably larger than 100 nm27–34. Thermally driven porosity formation in TMDs propagates through two mechanisms that operate in parallel: (1) reactive formation of gaseous H2X due to the thermally induced reaction
9
M (s) H 2 X ( g ) ), where X = S/Se and M is the metal) and with hydrogen ( MX (s) H 2 ( g )
(2) evaporation of chalcogen atoms from the TMD surface. Support for this mechanism is shown in Figure 2 comparing the morphology of thermally decomposed WSe2 at 950 °C in pure Ar (evaporative chalcogen removal only) to that in the presence of H2 (H2(5%)/Ar). The sample decomposed in H2/Ar is found to have more homogeneous porosity, while that in Ar has a discontinuous, bimodal porosity. Additionally, the final morphology length scale is smaller in H2/Ar than in Ar, 25 nm and 48 nm respectively, for the same decomposition conditions, 950 oC and 1 hr. This result indicates the reaction with H2 aids the fast removal of chalcogen atoms and consequently smaller initial porosity length scale. This behavior is also seen in classical electrochemical dealloying where higher rates of corrosion shift the balance between dissolution and surface diffusion, nucleating smaller length scales at early times22,53. Each of the two processes (evaporative and reactive) initiates the loss of chalcogen atoms through covalent bond severance at the gas-solid interface. This process is analogous to the oxidation and formation of a soluble cation in the electrochemical/chemical dealloying mechanism. As with traditional dealloying methodologies, the initiating event is the formation of a vacancy in the TMD lattice through the removal of a chalcogen atom. The exposure of low coordinated transition metal atoms on the surface induces a surface smoothening driven diffusion, mimicking spinodal decomposition of adatoms and vacancies, and formation of a porous architecture. The resulting morphology, as shown in Figure 1, resembles that of the bicontinuous nanoporosity observed in traditional dealloying systems16–25. 3.2
Effect of TMD precursor crystal structure
The precursor TMD crystal structure is found to have a significant effect on both the mechanism of porosity evolution and the final morphology/composition attained after processing. np-Ni and 10
np-Co are obtained through the thermal decomposition of pyrite-type TMDs54, Figure S1. These TMDs are defined by a long range crystal structure, covalent bonding throughout, where the metal atom is surrounded by six neighboring chalcogen atoms in a distorted octahedral orientation46,54,55. This more closely resembles the fcc structure of metal alloys in a traditional dealloying system. In contrast, layered TMDs possess a layered crystal structure where each MX2 layer (M = W, Re, Mo, Ti, Ta; X = Se, Se) is only connected to the neighboring layers through van der Waals attraction52,56. To understand the effect of pyrite-type vs. layered crystal structures, we compare the formation of np-Co and np-W. XRD, Figure 3(a), for the np-Co and np-W pre- and post-porosity evolution indicate the appearance of metallic phases of Co(002), CoSe(101), and CoSe(110) in np-Co and W(110) phase in np-W along with some residual TMD phases in the nanoporous material. The W(110) phase in np-W is also confirmed from the measured lattice parameter (0.2223 nm), Figure S2. Figure 3(b) shows the residual chalcogen contents in np-Co and np-W as a function of thermal decomposition time. For np-Co, the chalcogen content shows a steady decrease with time, approaching 25 at.% X (X = S, Se) after 9 hours. For np-W, the chalcogen content drops significantly in the first hour of decomposition. After 9 hours, np-W is found to be essentially pure W, with nearly all of the chalcogen removed. The time dependent loss of chalcogen content is confirmed through Raman spectroscopy (Figure S3). The change in A1g and E2g Raman modes57 of WSe2 are measured at different decomposition times. Both the out of plane A1g mode and in-plane E2g modes of WSe2 are visible up to 1 hour but disappear after 3 hours. Observation of the morphology and composition as it evolves with time and temperature for the pyrite-type, Co-based TMDs to np-Co shows a morphology and residual composition very similar to that of traditional dealloying, Figures 3, 6, and S1(b). The island formation and undercutting that define the early stages of dealloying, Figure 4(b), and
11
etch front propagation result in the formation of a core-shell structure, retaining a high fraction of the more reactive, less noble component within the interior of the ligaments4,5,19,20,22. In contrast to metal alloys and pyrite-type TMDs, only weak van der Waals forces hold MX2 layers together in layered TMDs (where M = W, Mo, Ti, Re etc.). This has two effects: (1) following chalcogen removal, there is no direct route, at the single atom scale, for the formation of a metal-metal bond between the transition metals in neighboring layers; (2) there is the potential for reactive gas intercalation between the MX2 layers. Regarding the latter, layered TMDs have been shown to be able to absorb/intercalate hydrogen, to the point where they have been evaluated as hydrogen storage materials58–61. This intercalation of reactive hydrogen facilitates chalcogen removal from both the exterior of the crystalline structure as well as the interior. The consequent result is the fast and complete reactive removal of the chalcogen element throughout the crystalline material, not simply limited to an etch front moving into the bulk of the crystal from a surface. Raman and EDS elemental analysis shows that thermal decomposition of the layered TMDs results in near complete chalcogen element removal, Figures 3 and S3. The interlayer chalcogen loss, to the point of full removal, along with the absence of any obvious mechanism for the formation of metal-metal bonds between the single atomic plane of metals in neighboring TMDs sheets suggests that the traditional dealloying mechanism is not operative here. Surface morphology at early times, 5 minutes of thermal decomposition, Figure 4, for the layered and pyrite-type TMDs show a distinct difference. For pyrite-type TMDs, classical mounding is observed and the initiation of the ligamentous structure is apparent. For the layered TMDs, however, the surface is decorated with small particles. With this observation, taken with the compositional data and the fact that H2 can intercalate between the layers, we can put forward a proposed mechanism for the formation of interconnected
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nanoporosity through the thermal decomposition of layered TMDs. As a consequence of the absence of metal-metal bonding between the MX2 layers and the intercalation of H2, the removal of chalcogen from the MX2 layers results in the coalescence of metal atoms into nanoscale clusters during thermal decomposition of the individual TMD layers. Occurring simultaneously internally and externally, the interconnected, bicontinuous nanoporosity then evolves through the aggregation and curvature driven smoothening of the aggregate structure as particles formed on adjoining layers coalesce. The observation of a small fraction of intra-ligament grain boundaries, Figure S4(a) and (b), for np-M formed from layered TMDs which are absent for np-M formed from the pyrite-type TMDs, Figure S4(c) and (d), supports this hypothesis. Figure 5 is a series of bright field TEM images taken during in-situ heating of WSe2 at 950 oC in an inert environment (10-5 Pa). This series of images shows the early stages of feature formation, initiating at the edges and propagating inward. Complete bicontinuous nanoporosity is attained after ~7 minutes. In-situ HRTEM video (Supporting Information), from which the still images in Figure 5 are obtained, shows an apparent uniform rate of etch front propagation and ligament formation progressing from the surface towards the bulk. This indicates that porosity evolution is likely interface reaction limited, which consists of chalcogen removal followed by transition metal diffusion and aggregation, and vapor phase chalcogen removal through both the pores and between the TMD layers is sufficiently fast. Further evidence for an interface limited process is presented in Figure 1(d) where there is an absence of a gradient of ligament size for porosity depths on the order of microns. The end result of the process is a sub-100 nm, uniform bicontinuous nanoporous metal with porosity depth on the order of microns (Figure 1(d)) for refractory metals (Re, Mo, W, Ta, Ti) which to this point has been very difficult to achieve otherwise. For these refractory metals, the degree of post-porosity evolution coarsening, even at
13
temperatures as high as 950 °C, is low due to the slow rate of surface diffusion. Surface diffusion rates of transition metals roughly scale inversely with the metal’s melting point62 . The np-W pore size remains below 100 nm even after 24 hr at 950 oC, Figure S5. 3.3
Rate controlling process: “Dissolution” vs. Diffusion
For nanoporosity formation through the removal of a sacrificial species, it is easy to think of the “corrosion” event as the defining process for both the rate of porosity evolution and the final morphology. Rate of removal of the more reactive chalcogen component of the TMDs is related to the composition of the TMD through the bond strength between the transition metal and chalcogen as well as the temperature dependence of the evaporation rate, related to the vapor pressure of the chalcogen and/or its reactivity with H2. In Figure 3(c) and (d)63, the experimentally observed minimum temperature required to initiate chalcogen loss and propagate nanoporosity is plotted versus both the melting point of the transition metal in the TMD as well as the formation energy of MS2 TMDs. A strong correlation is observed between the temperature required to initiate nanoporosity formation and the melting point of the transition metal within the TMD, Figure 3(c). In contrast, no significant correlation is observed with the formation energy of MX2. Experimentally we find that the identity of the chalcogen (X = S, Se) in layered TMDs, Figures S6, has no effect on the final morphology. Conversely, a modest effect of chalcogen identity is observed for the pyrite-type TMDs, in terms of average pore size, as a function of thermal decomposition temperature, Figure S7, and at a range of CoSxSe2-x compositions spanning from CoS2 to CoSe2, Figure 6(a)-(e). The M-Se bond is known to be stronger than the M-S bond46,64 and the vapor pressure of S is higher than that of Se, Figure 6(f). Therefore, MS2 should have a higher rate of thermal decomposition in comparison to MSe2 TMDs. The absence of any visible effect of chalcogen identity for the layered TMDs stems from 14
the low rate of surface diffusion for the high melting refractory metals from which they are composed. This low rate of surface reorganization dominates over any effects associated with chalcogen removal. The pyrite-type TMDs tested here are composed of transition metals with lower melting points, relative to the refractory metal containing layered TMDs, which corresponds to higher rates of surface diffusion at a given temperature. We argue that the observed effect of chalcogen identity for the pyrite-type TMDs, however modest, is due in part to the shift in the balance of the competition between “dissolution” and diffusion as a consequence of the higher rates of surface diffusion for the lower melting metals. The strong correlation between metal melting point and the minimum temperature required for nanoporosity formation in Figure 3(c) and the observed range in sensitivity to the chalcogen identity, from negligible for layered TMDs of higher melting transition metals to moderate for pyrite-type TMDs of lower melting transition metals, indicates that porosity formation and propagation is dominated by the surface diffusion of the remaining transition metal exposed through the removal of the chalcogen element. Activation energies for the thermally driven porosity evolution in CoS2 and CoSe2 are determined by plotting average ligament diameters versus temperature at early decomposition times (30 min.), Figure S8. The porosity evolution activation energies for CoS2 and CoSe2 are 99.96 kJ/mol (1.036 eV) and 104.02 kJ/mol (1.078 eV) respectively. Thermal evaporation of Zn from CoZn alloys was measured to have an activation barrier of 2.28 eV at low vacuum and 0.25 eV at high vacuum35. The difference in activation barrier was attributed to the shift in limiting processes. At low vacuum, the evaporation of Zn, speculated to be controlled by the bulk diffusion of Zn, is limiting. At high vacuum, the surface diffusion of Co is limiting35. The values obtained here lie within this range and suggest that porosity evolution is controlled by surface
15
diffusion of the transition metal. The absolute value of the activation barriers determined for CoS2 and CoSe2 thermal decomposition at atmospheric pressure, being closer to those determined at high vacuum conditions for vapor phase dealloying, is due to the reactive etching of the chalcogen elements in addition to simple evaporation. While the value of the activation barrier is suggestive of a surface diffusion limited process, there is a slight compositional dependence to the value, in line with data in Figure 6 and S7. This could be related to the slight difference in the rate of M-X bond breaking as the M-Se bond is stronger than the M-S bond46,64 and the vapor pressure of S is higher than Se, Figure 6(f). Overall, however, these results indicate a surface diffusion-controlled process, in agreement with other dealloying systems and methodologies, as well as the other data presented in this study. 4.
Conclusions In summary, we present a new, broadly applicable protocol for the synthesis of
morphologically controlled np-M with length scales below 100 nm, based on thermal decomposition of TMDs. With this new methodology, we have successfully synthesized a broad range of np-M where M = W, Re, Ta, Mo, Ti, Co, and Ni. Analysis of the impact of reactive gas, precursor crystal structure, and TMD composition gives insight into the mechanisms by which nanoporosity is evolved. Thermal decomposition offers a novel and universal pathway to synthesize a broad range of np-M from a diverse library of available TMDs.
Declaration of Competing Interest None.
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Figures
Figure 1: Scanning electron micrographs (SEM) of (a) np-W formed through the thermal decomposition of layered WSe2 at 950 °C under flowing H2(5%)/Ar (scale bar: 100 nm), (b) npMo: MoS2 precursor, layered, 800 °C, H2(5%)/Ar, 0.5 hrs. (scale bar: 100 nm), and (c) np-Re: ReS2 precursor, layered, 950 °C, H2(5%)/Ar, 0.5 hrs. (scale bar: 100 nm). (d) Focused ion beam (FIB) cross section of np-W (scale bar: 1 µm). (e) Schematic representation of the thermal decomposition process for the evolution of nanoporosity in TMDs.
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Figure 2: Demonstration of the effect of gaseous atmosphere (Ar vs. H2/Ar) on thermal decomposition of TMDs. (a) and (b) np-W from WSe2 after 1 hr. in Ar at 950 °C (scale bar: (a) 200 nm, (b) 2 µm). (c) and (d) np-W from WSe2 after 1 hr. in H2(5%)/Ar at 950 °C (scale bar: (a) 200 nm, (b) 1 µm). Images (a) and (c) are at the same magnification.
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Figure 3: (a) XRD spectra of CoSe2 (blue), np-Co (purple), WSe2 (black), and np-W (red). (b) Chalcogen content measured using SEM energy dispersive spectroscopy (EDS) as a function of thermal decomposition time for layered WSe2 (black) and pyrite-type CoSe2 (blue) crystalline precursor materials. (c) Minimum temperature required to initiate porosity formation through chalcogen removal and remaining metal surface diffusion in MX2 TMDs (X = S, Se; M = Ni, Co, Ti, Mo, Ta, Re, W) plotted versus the metal point of the pure transition metal. (d) Minimum temperature required to initiate porosity formation through chalcogen removal and remaining metal surface diffusion in MX2 TMDs plotted versus the formation energy of MS2. Red lines in (c) and (d) are linear fits to the data points.
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Figure 4: (a) SEM of the early stages of nanoporous morphology formation in WSe2 through thermal decomposition (5 min., 950 °C, H2(5%)/Ar), scale bar: 100 nm. (b) SEM of the early stages of nanoporous morphology formation in CoSe2 through thermal decomposition (5 min., 500 °C, H2(5%)/Ar), scale bar: 200 nm.
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Figure 5: In-situ heated stage TEM observing the time dependent morphology during thermal decomposition of WSe2 at 950 °C. (a) 0 minutes, (b) 4 minutes, (c) 5 minutes, (d) 7 minutes. Scale bars are 20 nm.
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Figure 6: SEMs of np-Co formed through the thermal decomposition (500 °C, 20 min, H2(5%)/Ar) of CoS(x)Se(2-x) with (a) x = 2, (b) x = 1.4, (c) x = 0.6, (d) x = 0. Scale bars are 1 µm. (e) Average pore size as a function of composition and thermal decomposition time; 20 min (black) and 30 min (red). (f) Vapor pressure as a function of temperature for S (blue) and Se (green).
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