*Article* **Application of Pulsed Laser Deposition in the Preparation of a Promising MoS***x***/WSe2/C(B) Photocathode for Photo-Assisted Electrochemical Hydrogen Evolution**

**Roman Romanov <sup>1</sup> , Vyacheslav Fominski 1,\* , Maxim Demin <sup>2</sup> , Dmitry Fominski <sup>1</sup> , Oxana Rubinkovskaya <sup>1</sup> , Sergey Novikov <sup>3</sup> , Valentin Volkov <sup>3</sup> and Natalia Doroshina <sup>3</sup>**


**Abstract:** We studied the possibility of using pulsed laser deposition (PLD) for the formation of a MoS*x*/WSe2 heterostructure on a dielectric substrate. The heterostructure can be employed for effective solar water splitting to produce hydrogen. The sapphire substrate with the conducting C(B) film (rear contact) helped increase the formation temperature of the WSe<sup>2</sup> film to obtain the film consisting of 2H-WSe<sup>2</sup> near-perfect nanocrystals. The WSe<sup>2</sup> film was obtained by off-axis PLD in Ar gas. The laser plume from a WSe<sup>2</sup> target was directed along the substrate surface. The preferential scattering of selenium on Ar molecules contributed to the effective saturation of the WSe<sup>2</sup> film with chalcogen. Nano-structural WSe<sup>2</sup> film were coated by reactive PLD with a nanofilm of catalytically active amorphous MoS*x*~4. It was established that the mutual arrangement of energy bands in the WSe<sup>2</sup> and MoS*x*~4 films facilitated the separation of electrons and holes at the interface and electrons moved to the catalytically active MoS*x*~4. The current density during light-assisted hydrogen evolution was above ~3 mA/cm<sup>2</sup> (at zero potential), whilst the onset potential reached 400 mV under irradiation with an intensity of 100 mW/cm<sup>2</sup> in an acidic solution. Factors that may affect the HER performance of MoS*x*~4/WSe2/C(B) structure are discussed.

**Keywords:** hydrogen evolution; pulsed leaser deposition; heterostructure; photoelectrocatalysis; semiconductors

#### **1. Introduction**

Transition metals chalcogenides have received considerable attention from scientists involved in the development of photoelectrochemical cells for producing hydrogen by solar water splitting [1–3]. These semiconducting materials have physio-chemical properties that enable their usage as both photo-active materials and hydrogen evolution electrocatalysts [4–6]. The good catalytic properties of metal chalcogenides (particularly amorphous molybdenum sulfides MoS*x*) allow the compounds to replace expensive platinum. Moreover, they can ensure high efficiency of photo-assisted hydrogen evolution when using silicon-based heterostructures (n+p-Si) [7–9]. The photoactivity of crystalline transition metal dichalcogenides is sufficient for creating photocathodes based on these materials. Efficient hydrogen evolution is usually achieved by using an expensive (Pt/Ru) cocatalyst [10,11].

It is essential to coalesce the useful semiconductive and catalytic properties of transition metal chalcogenides in creating hybrid or heterostructures. These structures consist entirely of thin-film metal chalcogenides that have been selected for their structure and

**Citation:** Romanov, R.; Fominski, V.; Demin, M.; Fominski, D.; Rubinkovskaya, O.; Novikov, S.; Volkov, V.; Doroshina, N. Application of Pulsed Laser Deposition in the Preparation of a Promising MoS*x*/WSe2/C(B) Photocathode for Photo-Assisted Electrochemical Hydrogen Evolution. *Nanomaterials* **2021**, *11*, 1461. https://doi.org/ 10.3390/nano11061461

Academic Editors: Antonio Riveiro, Jesús del Val and Mohamed Boutinguiza

Received: 26 April 2021 Accepted: 30 May 2021 Published: 31 May 2021

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**Copyright:** © 2021 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

chemical composition. These compounds also act as photo-assistant agents in water splitting for hydrogen production [12].

Tungsten diselenide is a promising photo-active transition metal dichalcogenide [12–15]. Crystalline WSe<sup>2</sup> is a p-type semiconductor with a small band gap (~1.1 eV). If the conductivity is sufficiently high, this compound can be combined with catalytically active n-type metal chalcogenides to form photocathodes. The central requirement for WSe<sup>2</sup> films used in such photocathodes arises from the need to obtain a nearly perfect structure with minimum defects, including edge states. The recombination rate of nonequilibrium carriers (electrons and holes), which form during irradiation, decreases in the process. However, nanostructured WSe<sup>2</sup> films can have a greater area for hydrogen evolution, while edge states can be passivated by a co-catalyst [10,13]. WSe<sup>2</sup> films with a nearly perfect crystal lattice are usually obtained by chemical synthesis (from the vapour phase or in a special solution) or by the selenization of thin-film precursors (for instance, [16,17]). These techniques have both advantages and disadvantages. Thus, finding alternative techniques to obtain thin WSe<sup>2</sup> film with targeted properties remains a challenge.

In its traditional on-axis configuration, pulsed layer deposition (PLD) makes it possible to create WSe<sup>2</sup> with a crystal structure and good catalytic properties [18–21]. However, when using the on-axis PLD to obtain WSe<sup>2</sup> films, problems were revealed with obtaining a stoichiometric composition with a perfect chemical state of atoms in the film. This situation can be attributed to several factors, including the preferential sputtering of selenium on a growing WSe<sup>2</sup> when exposed to laser-plasma and the propensity of selenium to form pure Se nanoparticles at room temperature or to be desorbed at higher substrate temperatures [20,22–24]. Submicron- and nanoscale particles of metal W can be introduced into the film [25]. These particles form upon laser irradiation of a WSe<sup>2</sup> target.

In the case of on-axis PLD geometry, the substrate is placed normally to the axis of the laser plume expansion. Fominski et al. [26] have established that, under some on-axis PLD regimes, using a buffer gas makes it possible to decrease the efficiency of the preferential self-sputtering of chalcogen atoms. However, this technique suffers from considerable limitations: under some regimes of ablation of a dichalcogenide metal target, a laser plume may form and localize in a narrow solid angle. It has been revealed that chalcogen atoms can move to the plume periphery, while a buffer gas cannot preclude the metallization of the centre of the film deposition area [27].

Experiential studies of amorphous molybdenum sulfide (a-MoS*x*) catalytic film creation have shown that, during off-axis PLD, the buffer gas has the conditions necessary for the effective saturation of films with chalcogen atoms [28]. These conditions are a result of the difference between the S and Mo atomic masses. During off-axis PLD, the substrate is placed parallel to the plume expansion. The film grows chiefly through the deposition of atoms that have collided with buffer gas molecules and changed their direction. The same effect can be achieved during the off-axis PLD of WSe<sup>2</sup> films due to the substantial difference between the W and Se atomic masses. One of the goals of this study was to test whether the off-axis PLD technique could be applied to a WSe<sup>2</sup> target to obtain WSe<sup>2</sup> nanocrystal films that are nearly perfect in terms of structure and chemical state.

Films based on a-MoS*x*-nanomaterial have a high electrocatalytic activity during hydrogen evolution reaction (HER) [7,13,29–31]. The most common technique to obtain such films is a chemical synthesis or chemical deposition in the solutions of special precursors. The applicability of a laser-based technique to create amorphous a-MoS*<sup>x</sup>* films with good electrocatalytic properties has been described in the literature [28,32–35]. Fominski et al. demonstrate that the PLD technique is associated with the highest catalytic activity in MoS*x*~4 films because they contain Mo3S12/Mo3S13 clusters [35]. It is suggested that reactive PLD (RPLD) in H2S gas be used to create thin homogenous a-MoS*<sup>x</sup>* films with a high content of catalytically active states of sulfur [34]. Performing on-axis RPLD from a Mo target prevents the formation of a substantial number of particles of various sizes during ablation. It also ensures a relatively conformal coating of a rough surface, which is typical of catalysts. This beneficial effect is possible because deposition is carried out using

a flux of Mo atoms, which scatter at different angles once collided with H2S molecules. Given the difficulty of predicting the conductivity type in a-MoS*<sup>x</sup>* films with increased S content, it was necessary to establish whether the mutual arrangement of energy bands in the MoS*x*~4/WSe<sup>2</sup> heterostructure is optimal for the effective separation of electron-hole pairs when light irradiated during the photo-assisted HER.

It is widely accepted that the material of the rear contact to the semiconductor can have a pronounced effect on the current transport in the semiconductor photovoltaic structure and probably also the photoelectrocatalyst [36]. We investigated boron-doped carbon films as rear contacts. A preliminary study showed that the introduction of boron atoms could produce p-type conductivity of C(B) films [37]. Films with good conductivity and mechanical strength were created by PLD from a mixed boron/graphite target. The substrate of the heterostructure was a sapphire plate. However, sapphire can be replaced by a cheaper material—glass, quartz, etc.

Our study aimed to form a multi-layered MoS*x*~4/WSe2/C(B) by PLD. The structure had to contain thin-film nanomaterials with properties sufficient for an effective photoassisted HER in an acidic solution. When selecting PLD conditions for obtaining these nanomaterials, we used the results of a preliminary investigation of each nanomaterial. After the heterostructure have been assembled (i.e., after layer-by-layer nanomaterial deposition), chosen PLD conditions may prove to be non-optimal for efficient photoelectrocatalysis of hydrogen evolution. Nonetheless, the findings made it possible to produce recommendations on how laser-based processes may be improved and the structure and composition of selected nanomaterials modified.

#### **2. Materials and Methods**

#### *2.1. Experimental Methods for On-Axis and Off-Axis PLD of Functional Nanolayers for MoSx~4/WSe2/C(B) Heterostructure Formation*

Figure 1 shows the mutual arrangement of the target and the substrate when using on-axis and off-axis PLD for the formation of a MoS*x*~4/WSe2/C(B) heterostructure. The on-axis PLD configuration was used to deposit a C(B) film. The target, which consisted of a carbon (soot) and boron powder mixture in the proportion C/B~6, was concurrently ablated. For more detail on target manufacturing and the selection of laser ablation, see [37]. A Solar LQ529 laser (Minsk, Belarus) was used to ablate the target. The pulse duration and energy were 10 ns and 100 mJ respectively. The pulse repetition rate was 20 Hz. The energy density on the surface of the C(B) target was 9 J/cm<sup>2</sup> . The substrate was placed to the laser plume axis, 3 cm away from the target, and heated to 500 ◦C. The deposition was performed in a vacuum at a residual pressure of 5<sup>×</sup> <sup>10</sup>−<sup>4</sup> Pa. The deposition period of C(B) films was 10 min. The film thickness did not exceed 150 nm.

**Figure 1.** A schematic of the PLD technique employed to form functional layers in a MoS*x*~4/WSe2/C(B) heterostructure on a sapphire substrate. Comments are given in the text.

A WSe<sup>2</sup> film was deposited by off-axis PLD on the surface of the substrate coated with a C(B) film. The substrate was rotated 90◦ and placed along the laser plume axis 2 cm away from the WSe<sup>2</sup> target, which was manufactured by cold pressing of WSe<sup>2</sup> powder [22]. During the WSe<sup>2</sup> target ablation, the laser fluence was reduced to 4 J/cm<sup>2</sup> . The substrate temperature was 700 ◦C. WSe<sup>2</sup> film deposition was performed in an Ar + 5% H<sup>2</sup> mixture at a pressure of 15 Pa. The gas mixture was introduced into a chamber that had been evacuated to a residual gas pressure of 5 <sup>×</sup> <sup>10</sup>−<sup>4</sup> Pa or less. The deposition time of a WSe<sup>2</sup> film of a thickness of ~200 nm was 20 min.

After the formation of WSe<sup>2</sup> films, the sample was allowed to cool down to room temperature. Then it was rotated 90◦ for subsequent MoS*<sup>x</sup>* film deposition by PLD in the reaction gas. The Ar + 5% H<sup>2</sup> gas mixture was pumped out with the help of a turbo-molecular pump, and H2S gas was introduced into the chamber until the pressure reached ~26 Pa. The Mo target ablation was carried out using 100 mJ pulses. The MoS*<sup>x</sup>* film deposition time was set at 6 min. The thickness of a deposited MoS*<sup>x</sup>* film on a smooth substrate did not exceed 20 nm. The choice of the Mo ablation conditions and the H2S pressure was motivated by the results of preliminary investigation of RPLD of MoS*<sup>x</sup>* films [38]. Under the chosen conditions of on-axis RPLD, the expected ratio was = S/Mo~4.

#### *2.2. Structural, Chemical, Electrical, Optical, and Photoelectrochemical Characterization Techniques*

In this study, WSe<sup>2</sup> films were produced by off-axis PLD for the first time, and thus they require further examination. Yet, C(B) and MoS*<sup>x</sup>* films obtained by on-axis (R)PLD have been studied extensively. We also discussed their structural and chemical properties in several publications. Therefore, in this article, we will focus on the information that will give a comprehensive picture of the components (layers) of the MoS*x*~4/WSe2/ C(B) heterostructure.

The surface morphologies of the prepared films and heterostructures were examined by scanning electron microscopy (SEM, Tescan LYRA 3, Brno, Czech Republic). Using this microscope, the surface distribution of elements was studied by energy dispersive X-ray spectroscopy (EDS). The structure of the films was investigated by micro-Raman spectroscopy (MRS, Horiba, Kyoto, Japan), using a 632.8-nm (He-Ne) laser. The crosssection of the laser beam was <1 µm. To explore the structural features of WSe<sup>2</sup> films obtained by off-axis PLD, the films were separated from the substrate and transferred onto metal grids to study by high-resolution transmission electron microscopy (HRTEM) and selected area electron diffraction (SAED) with the help of a JEM-2100, JEOL microscope (Toyo, Japan).

Band gaps (*E*g) in the prepared films were measured optically by processing absorption spectra. To this end, a Tauc plot was constructed to describe the dependence between (*αhν*) 1/r and (*hν*), where *α* is the absorption coefficient, *hν* is the photon energy, and *r* is a parameter that is taken to be 2 for indirect transitions. The optical absorption and transmission spectra were measured using an Agilent Technologies Cary Series UV-Vis-NIR spectrophotometer. Special samples were manufactured to explore the optical properties of WSe<sup>2</sup> and MoS*<sup>x</sup>* films. In these samples, WSe<sup>2</sup> and MoS*<sup>x</sup>* films were deposited on transparent sapphire substrates in the selected conditions.

The chemical states of WSe<sup>2</sup> and MoS*<sup>x</sup>* films were studied by XPS. XPS spectra were obtained using a Theta Probe Thermo Fisher Scientific spectrometer with a monochromatic Al Kα X-ray source (*hν* = 1486.7 eV) and a 400 µm X-ray spot. The spectrometer energy scale was calibrated using Au4f7/2 core level lines located at a binding energy of 84.0 eV. The Advantage Data Spectrum Processing program was used for deconvolution of the experimental XPS spectra. The Shirley background is an approximation method that was used for determining the background under an XPS peak. The peaks were fitted by symmetric convolution of Gaussian and Lorentzian functions. The ratios of atomic concentrations of elements (*x* = S/Mo) were calculated considering the intensities of Mo 3d and S 2p peaks and the corresponding Scofield's Relative Sensitivity Factor.

The thickness of quite thick MoS<sup>x</sup> films (thickness is ≥100 nm) was measured by SEM. For this, the Si substrate with the deposited thick MoS*<sup>x</sup>* film was cleaved and the vertical cross section was investigated by SEM. These measurements made it possible to estimate the deposition rate of the MoS*<sup>x</sup>* films during reactive PLD. The deposition

rate was used to determine the time for preparation of a very thin MoS*<sup>x</sup>* film (thickness is ~3 nm). The thickness of this thin film was then estimated from the results of XPS studies for MoS*x*/WSe<sup>2</sup> heterostructure. The thickness of thin MoS*<sup>x</sup>* film was estimated as quite adequate if the XPS spectra of both films (MoS*<sup>x</sup>* and WSe2) could be detected at the same time.

The XPS measurements were used to determine the mutual arrangement of valence bands (VB) in the semiconductor heterostructures. The employed technique is widely used to study the band structure in heterojunction the formation of which can cause a change in the energy distribution of electrons [35,39,40]. The leading edge of the valence band spectrum was approximated by a linear function using the least-square fit of the leading edge of the VB spectra. The position of the valence band maximum (VBM) was determined as the intersection of the approximating linear function and the baseline. Determining the shift between the core levels of semiconductors in the heterojunction made it possible to calculate the valence band offset (VBO).

To calculate the VBO in a MoS *<sup>x</sup>*/WSe<sup>2</sup> heterostructure, a series of measurements was performed. Firstly, the XPS spectra of the Mo3d and W4f core levels were measured along with the spectra of the valence bands of quite thick MoS*<sup>x</sup>* and WSe<sup>2</sup> films. Secondly, the spectra of the Mo3d and W4f core levels were measured for a MoS *<sup>x</sup>*~4/WSe<sup>2</sup> heterostructure, in which the thickness of the upper layer (MoS) did not exceed 3 nm. Thirdly, the VBO value for heterojunctions was calculated based on the formula:

$$\text{VBO} = (\text{E}\_{\text{Mo3d5}/2} - \text{E}\_{\text{W4f7}/2})\_{\text{interface}} + (\text{E}\_{\text{W4f7}/2} - \text{VBM}\_{\text{W}})\_{\text{bulk}} - (\text{E}\_{\text{Mo3d5}/2} - \text{VBM}\_{\text{Mo}})\_{\text{bulk},1}$$

where *VBM<sup>W</sup>* and *VBMMo* are the energies of the upper edge of the valence band for WSe<sup>2</sup> and MoS, respectively. 'Interface' stands for spectra for heterojunctions, and 'bulk' for the spectra of thicker films on C(B)/Al2O<sup>3</sup> substrates.

The work function (ϕ) needed to withdraw an electron from a WSe<sup>2</sup> film was calculated using the formula ϕ = *hν* − *E CutOff* + *EF*, where *E CutOff* is the secondary electron cutoff, *E*<sup>F</sup> is the Fermi level if these magnitudes are considered on a kinetic energy scale. The Fermi level was determined based on an analysis of the energy spectrum of the valence band. To enable XPS investigation, the samples were created on a conducting C(B) film. This way, charge storage was prevented in the sample. The zero-value point of the binding energy scale corresponded to the Fermi level. In this case, the *E CutOff* value marked on the kinetic energy scale coincides with ϕ.

The electrical properties of C(B) films on sapphire substrates were studied by a fourcontact method in the van der Pauw geometry; Hall-effect measurements were performed at room temperature. A magnetic field, varying from 0 to 1 T, was used for Hall-effect characterization. Metallic contacts to the sample with a circular mesa were formed from an InSn alloy, and the linearity of the volt–ampere characteristics of all contacts was monitored. During the measurements, the direction of the current was switched to eliminate the effects of thermoelectric power. The resistivity was calculated by averaging the values from all pairs of contacts.

To study the photoelectrocatalytical properties of MoS*x*/WSe2/C(B)/Al2O<sup>3</sup> samples, we irradiated these samples with 100 W Xe lamps in an 0.5 M H2SO<sup>4</sup> aqueous solution. The light intensity was 100 mW/cm<sup>2</sup> . A three-electrode configuration was used to determine the photo-assisted current in an electric circuit with modified cathodes. The polarization curves were measured using linear sweep voltammetry (LSV) with a change in the applied potential from −100 to 400 mV and a scan rate of 2 mV/s. When measuring LSV curves and the time evolution of the photocurrent, the light source was turned on and off. For chronoamperometry measurements, the potential of the photocathode was 0 V (relative to the reversible hydrogen electrode, RHE).

#### **3. Results**

#### *3.1. On-Axis PLD of C(B) Films*

Figure 2a,b show the morphology of a C(B) film formed on sapphire by traditional PLD. Detached rounded particles were observed on the smooth surface of C(B) films. The particle size ranged from 0.1 to 0.5 µm. This morphology is attributed to the deposition of B-rich particles [37]. The Raman spectrum of a C(B) film has two broad peaks at 1343 and 1545 cm−<sup>1</sup> (Figure 2c), which correspond to the peaks marked D and G. The peaks are shifted to the lower wavenumber relative to the peaks associated with graphite (1360 cm−<sup>1</sup> and 1580 cm−<sup>1</sup> respectively). This shift meant that the films had a graphite-like local packing, which contained B atoms and some C atoms with sp<sup>3</sup> -bonding in sp<sup>2</sup> -matrix. A more detailed analysis of the Raman spectra of C(B) films obtained by PLD can be found in [37,41].

**Figure 2.** SEM images obtained at (**a**) normal and (**b**) 45◦ angles to the surface; (**c**) Raman spectrum of a C(B) film prepared on sapphire substrate by on-axis PLD in a vacuum.

The C(B) films had a specific resistance of ~1.5 mΩ·cm and p-type conductivity. At room temperature, the carrier concentration and mobility were 4.4 <sup>×</sup> <sup>10</sup><sup>19</sup> cm−<sup>2</sup> and 180 cm2/V·s, respectively. The low resistance to the current flow in C(B) films enabled their use as a rear contact to the MoS*x*~4/WSe<sup>2</sup> heterostructure, and p-type conductivity made it possible for holes formed upon illumination to move from the WSe<sup>2</sup> film to the external electric circuit.

#### *3.2. Off-Axis PLD of WSe<sup>2</sup> Films*

Figure 3 shows the morphology of the WSe<sup>2</sup> film obtained by off-axis PLD on the surface of C(B) film. The WSe<sup>2</sup> film covers the surface of the C(B) film with a continuous layer and the WSe<sup>2</sup> film had a nanocrystal structure consisted of petal-like crystals with random orientation relative to the film surface. The linear sizes of WSe<sup>2</sup> crystals reached 1 µm, whereas the thickness of the nanopetals did not exceed 50 nm.

**Figure 3.** (**a**,**b**) SEM images of the surface of the WSe<sup>2</sup> film (two magnifications) obtained by off-axis PLD on the surface of C(B)/Al2O<sup>3</sup> sample.

A structural investigation by MRS and TEM/MD techniques demonstrated that the WSe<sup>2</sup> film had a crystal structure. An MRS spectrum (Figure 4a) only had peaks characteristic of the 2H-WSe<sup>2</sup> phase. The peaks associated with the vibrational modes E2g <sup>1</sup> and A1g coincided because the shift between them was approximately 3 cm−<sup>1</sup> [42,43]. The narrow half-height width of the peak (3 cm−<sup>1</sup> ) points to the suitable quality of the crystal structure. A high-resolution TEM and SAED analysis of a single WSe<sup>2</sup> petal showed that it consisted of several nanocrystals with a hexagonal lattice of the 2H-WSe<sup>2</sup> phase (Figure 4b). The nanocrystal size was ~10 nm. Although the nanocrystals were oriented randomly relative to the *c*-axis, the basal plane of all the nanocrystals was parallel to the petal surface.

**Figure 4.** (**a**) Raman spectra and (**b**) HRTEM and SAED patterns of the WSe<sup>2</sup> film obtained by off-axis PLD.

Figure 5 shows part of the XPS spectra for the surface of WSe<sup>2</sup> film deposited on the surface of C(B) film. The W4f spectrum was well described by a doublet in which the W4f7/2 and W4f5/2 peaks had binding energies of 32.24 and 34.47 eV, respectively, which are characteristic of WSe2. The Se3d spectrum was described by a doublet whose Se3d5/2 and Se3d3/2 peaks were at 54.50 and 55.37 eV, respectively. The XPS spectra indicated effective chemical interaction between Se and W during off-axis PLD [13,43,44].

**Figure 5.** Core level XPS W4f and Se3d spectra of the WSe<sup>2</sup> film obtained by off-axis PLD.

An analysis of the energy spectrum of secondary electrons and the valence band showed that the work function for WSe<sup>2</sup> electrons was 4.9 eV (Figure 6a). The Fermi level was close to the bottom of the band gap 0.25 eV away from the upper edge of the valence band (Figure 6b). An investigation of the WSe<sup>2</sup> film optical properties demonstrated that the film had an absorption spectrum characteristic of WSe2; the band gap width was 1.4 eV (Figure 7). A study of the band structure of the WSe<sup>2</sup> film proved that it had p-type conductivity typical of this compound [45,46].

**Figure 6.** (**a**) XPS spectrum of secondary electrons cutoff and (**b**) evaluation of the valence band edge position for a WSe<sup>2</sup> film obtained by off-axis PLD on the surface of C(B) film.

**Figure 7.** (**a**) Optical absorption spectra and (**b**) Tauc plots for the WSe<sup>2</sup> film on the sapphire substrate. The A and B peaks are explained by excitonic absorption.

#### *3.3. On-Axis Reactive PLD of MoSx~4 Film*

Figure 8 shows an SEM image of the surface of WSe2/C(B)/Al2O<sup>3</sup> sample after MoS*<sup>x</sup>* film deposition by on-axis reactive PLD. MoS*<sup>x</sup>* film deposition did not cause a substantial change in the morphology of the sample surface. The principal difference between SEM images of the WSe<sup>2</sup> film before (Figure 3) and after MoS*<sup>x</sup>* film deposition (Figure 8) was that the sides of the WSe<sup>2</sup> nanocrystal petals lost their sharpness when coated by a thin MoS film, which is a porous structure. Mapping element distribution in the sample surface suggested that the MoS*<sup>x</sup>* film had a sufficiently homogeneous distribution over the sample surface (Figure 9). During PLD, the collision of Mo atoms with H2S molecules ensured their scattering at different angles. As a result, a MoS*<sup>x</sup>* film could be formed even on those WSe<sup>2</sup> nanopetals that were oriented perpendicular to the surface of the substrate.

**Figure 8.** (**a**,**b**) SEM images (two magnifications) of the MoS*x*/WSe2/C(B)/Al2O<sup>3</sup> sample.

**Figure 9.** SEM image (top gray) and EDS maps (colored) of element distribution on the surface of the MoSx/WSe2/C(B)/Al2O<sup>3</sup> sample. Intensity of different colors indicates where the corresponding elements (Se, W, S, Mo, B, and C) are most abundant. The presence of submicron rounded particles is explained by B-rich particle deposition during C(B) film formation.

Figure 10 shows the results of XPS investigation of a quite thick MoS*<sup>x</sup>* film obtained by on-axis RPLD. An analysis of the chemical state of elements showed (Figure 10a,b) that core level XPS Mo3d spectrum was well described by a doublet corresponding to the Mo4+

state. The bonding energy of the peak Mo3d5/2 was 229.24 V, accounted for by chemical bonds with S atoms [26,34]. Molybdenum oxides (Mo6+) or metallic <sup>0</sup> were not observed. The Mo3d5/2 peak partially overlapped with the S2s peak. The S2s peak consisted of singular peaks whose position correlated with that of doublets in the S2p spectrum. When analyzing the S2p peak, we used the traditional approach, i.e., we identified the states of sulfur with high and low binding energy (HBE and LBE, respectively) [34,35]. The LBE doublet was associated with single S2- atoms (in MoS2-like clusters) and a terminal (S<sup>2</sup> <sup>2</sup>−)tr ligand (in Mo3S13/Mo3S<sup>12</sup> clusters). The doublet had S2p3/2 and S2p1/2 peaks with binding energies of 162.04 and 163.35 eV, respectively. The HBE doublet had S2p3/2 and S2p1/2 peaks, whose binding energies were 163.28 and 164.50 eV. This doublet is usually attributed to apical S2- and bridging (S<sup>2</sup> <sup>2</sup>−)br ligands in Mo3S13/Mo3S<sup>12</sup> clusters. An XPS studies-based calculation of S/Mo atomic concentration ratios for this film confirmed that *x*~4.0. Measuring the valence band spectrum showed that the Fermi level was 0.4 eV away from the bottom of the band gap (Figure 10c).

**Figure 10.** XPS spectra of (**a**,**b**) core level Mo3d and S2p and (**c**) the valence band of a relatively thick MoS*x* film obtained by on-axis RPLD.

Figure 11a shows the spectrum of optical absorption for the MoS*x*~4 film. Figure 11b demonstrates a Tauc plot calculated for that spectrum. The optical properties of MoS*x*~4 films are very similar to those of WSe2. This similarity sets a limit on the thickness of the MoS*<sup>x</sup>* film in a MoS*x*/WSe<sup>2</sup> heterostructure since both films absorbed light most efficiently at wavelengths below 500 nm. The width of the band gap in a MoS*<sup>x</sup>* film was 1.55 eV. A Fermi level in the lower part of the band gap indicated p-type conductivity in the MoS*x*~4 film.

**Figure 11.** (**a**) Optical absorption spectra and (**b**) Tauc plots for the MoS*x*~4 film deposited on sapphire substrate.

The local packing of atoms in the MoS*x*~4 film was investigated by MRS. It can be seen in Figure 12 that the Raman spectrum of the film consists of a set of broadened

strips, whose position correlates well with that of the bands in the Raman spectrum of a catalytic molybdenum sulfide film obtained by chemical synthesis in a solution [13] and by reactive magnetron sputtering [47]. The spectrum had two clear broadened peaks at ~525 and ~550 cm−<sup>1</sup> , which were accounted for by the vibrational modes *ν*(S-S)tr and ν(S-S)br respectively in Mo3S13/Mo3S<sup>12</sup> clusters. The peak at ~450 cm−<sup>1</sup> is explained by the vibrations of apical S in Mo3S<sup>13</sup> clusters. A broad band in the range 250–400 cm−<sup>1</sup> is characteristic of an amorphous featureless structure of MoS*x*. Thus, the selected regime of on-axis RPLD made it possible to obtain thin layers of an amorphous molybdenum sulfide containing Mo3S13/Mo3S<sup>12</sup> clusters on the surface of a nanostructured WSe<sup>2</sup> film. The high electrocatalytic activity of such an amorphous molybdenum sulfide could contribute to a photo-assisted HER if the flux of nonequilibrium carriers (electrons) through the interface with WSe<sup>2</sup> was sufficient.

**Figure 12.** Raman spectra of MoS*x*~4 film obtained by on-axis RPLD on a sapphire substrate.

## *3.4. Photoelectrocatalytic Properties of the MoSx~4/WSe2/C(B)/Al2O<sup>3</sup> Cathode*

Figure 13 shows the results of an investigation of photoelectocatalytic properties of various heterostructure based on laser-deposited MoS*x*~4, WSe2, and C(B) films on a sapphire substrate. During photo-assisted HER, the MoS*x*~4/WSe2/C(B) materials combination had the most suitable properties (Figure 13a). A luminous flux caused the photo-current density to increase to ~3 A/cm<sup>2</sup> at a voltage of 0 V(RHE). The photocurrent magnitude was superimposed with the relatively high dark current raised due to transient effects [43]. The onset potential reached 400 mV (RHE). The heterojunction between MoS*x*~4 and WSe<sup>2</sup> films was largely responsible for an efficient photo-assisted HER in this photocathode. Cathodes with a single semiconductive layer (MoS*x*~4 or WSe2) on the C(B) layer were associated with very low efficiency of photo-assisted HER (Figure 13b).

In the study of the temporal stability of the MoS*x*~4/WSe2/C(B)/Al2O<sup>3</sup> photocathode, the current density was found to rapidly decrease by 20% in 20 min under chopped illumination. After a period of decline, the current density remained relatively stable for two hours. Longer tests of the temporal stability of this photocathode were not performed.

Table 1 contains collected data for comparison of the main parameters of metal chalcogenide-based photocathodes that characterize their performance in photo-assisted HER. It can be seen that the MoS*x*~4/WSe2/C(B)/Al2O<sup>3</sup> photocathode created by pulsed laser deposition is not inferior in general in photo-assisted HER to the performance of photocathodes which were prepared by the methods of wet/dry chemical synthesis, exfoliation, spin coating, etc. Next, we will discuss the factors that should be overcome to enhance the photo-assisted HER efficiency of the MoS*x*~4/WSe2/C(B)/Al2O<sup>3</sup> photocathode.

**Figure 13.** (**a**) Chopped LSV curve for the MoS*x*~4/WSe2/C(B)/Al2O<sup>3</sup> photocathode in 0.5 M H2SO<sup>4</sup> upon illumination; (**b**) Chopped photocurrent density versus time for MoS*x*~4/WSe2/C(B)/Al2O<sup>3</sup> (curve *1*), MoSx~4/C(B)/Al2O<sup>3</sup> (curve *2*) and WSe2/C(B)/Al2O<sup>3</sup> (curve *3*) photocathodes at 0 V (RHE) in 0.5 M H2SO<sup>4</sup> upon illumination.

**Table 1.** Comparison of photo-assisted HER performances for metal chalcogenide-based photocathodes with heterojunction structure.


\* 1 the amorphous solid–liquid–crystalline solid process with Pd promoter. \*<sup>2</sup> selective dip coating impregnation.

#### **4. Discussion**

XPS studies of MoS*x*~4 and WSe<sup>2</sup> layers obtained by PLD showed that they had p-type conductivity. Such a combination of the electrophysical properties of contacting semiconductors creates a situation when the efficiency of photo-assisted HER processes largely depends on the structure of energy bands at the MoS*x*~4 /WSe<sup>2</sup> interface. Figure 14 shows band alignment at the interface which was determined through a comprehensive study of the films by XPS and optical methods. The conductive band offset (CBO) value was calculated using the formula:

$$\text{CBO} = \text{VBO} + E\_{\S}(\text{WSe}\_2) - E\_{\S}(\text{MoS}\_{x-4}).$$

**Figure 14.** Band alignment diagram for the MoS *<sup>x</sup>*~4/WSe<sup>2</sup> heterojunction system obtained by PLD.

The obtained CBvalue equaled 0.1 eV. Thus, the band alignment was of type II, which is associated with the most efficient separation of photo-generated electron-hole pairs. In this case, electrons will move into the MoS*<sup>x</sup>* layer from the WSe<sup>2</sup> and participate in the hydrogen evolution reaction, whilst holes will move from MoS*<sup>x</sup>* into the WSe<sup>2</sup> layer. From WSe2, holes will migrate into the C(B) rear contact and further into the external electric circuit.

An additional study of MoS*x*~4/WSe2/C(B)/Al2O<sup>3</sup> samples by electrochemical impedance spectroscopy (EIS) demonstrated that the C(B) contact layer did not ensure a sufficiently low resistance to the flow of current. The value of equivalent series resistance (*R*s), which was extracted from EIS data, achieved 30 Ω. When a glassy carbon conducting substrate was used to create a MoS*x*~4/WSe2/GC photocathode, *R*<sup>s</sup> did not exceed 4 Ω, and the density of the photo-assisted HER current increased. Therefore, to increase the efficiency of a photo-assisted HER when using a MoS *<sup>x</sup>*~4/WSe<sup>2</sup> heterojunction system, it is recommended to choose a rear contact with an electrical resistance lower than that of the C(B) film. Further work may focus on the effect of the B concentration on the electrical properties of such films.

The analysis of the optical characteristics of the MoS*x*~4 and WSe<sup>2</sup> films showed that these films absorb light rather efficiently. These nanomaterials are potentially active catalysts for the hydrogen evolution reaction. However, these factors did not provide effective photo-assisted HER in the MoS*x*~4/C(B)/Al2O<sup>3</sup> and WSe2/C(B)/Al2O<sup>3</sup> samples. Additional experiments with thicker MoS*x*~4 and WSe<sup>2</sup> films did not reveal significant changes in the efficiency of photo-assisted HER. This indicated that after the generation of electron-hole pairs under a light flux, electrons and holes could rapidly recombine in the bulk of the films. The formation of a heterojunction turned out to be the most important factor contributing to an increase in the photocurrent. At the interface of the MoS*x*~4 and WSe<sup>2</sup> films, not only the processes of separation of nonequilibrium electrons and holes due to the specificity of the energy bands alignment could occur, but also recombination processes can be expected. The recombination processes will facilitate to photo-assisted HER if electrons from WSe<sup>2</sup> and holes from MoS*x*~4 actively participated in the recombination process (Z-schema) [35]. However, one cannot exclude the recombination at this interface of electrons and holes generated by the light flux in the WSe<sup>2</sup> film. In addition, the small size of the crystalline domains in the WSe<sup>2</sup> film and their random orientation resulted in a high density of edge states. This should lead to a decrease in the efficiency of charge separation since such edge states serve as recombination centers [13]. Insufficiently large values of CBO and VBO for the MoS*x*~4/WSe<sup>2</sup> heterojunction could also be the reason limiting the efficiency of photo-assisted HER in our samples.

Another factor that could reduce the efficiency of a photo-assisted HER with a MoS*x*~4/WSe2/C(B)/Al2O<sup>3</sup> photocathode is modification of the MoS*x*~4/WSe<sup>2</sup> interface under the influence of hydrogen sulfide activated by laser-induced plasma. Figure 15 shows W4f and Mo3d XPS spectra measured for a very thin MoS*<sup>x</sup>* film formed by on-axis RPLD on the surface of the WSe<sup>2</sup> layer. A comparison of these spectra with those of pristine WSe<sup>2</sup> and MoS*x*~4 (Figures 5 and 10) demonstrated that the chemical state of W has practically not changed after the deposition of MoS*<sup>x</sup>* film. The Mo3d spectrum shifted by 0.37 eV towards greater bonding energies, whilst the S2s spectrum increased in intensity. These changes indicated that, at the initial stage of the MoS*<sup>x</sup>* film growth, sulfur could be effectively deposited on the WSe<sup>2</sup> as a result of H2S molecules interacting with the WSe<sup>2</sup> surface. The plasma that formed in H2S during the ablation of the target could activate the process. The introduction of S atoms into the WSe<sup>2</sup> crystal lattice can distort the latter and thus cause the formation of new energy levels in the WSe<sup>2</sup> band gap. At the same time, energy bands will bend in the contact area. Band bending may cause a bonding energy shift for the Mo3d5/2 peak and increase the width at the half maximum of the peak from 1.4 to 1.8 eV. At these energy levels, effective recombination of electrons and holes formed upon illumination may occur.

**Figure 15.** XPS W4f and Mo3d (overlapped with S2s) spectra for a very thin MoS*<sup>x</sup>* film obtained by on-axis reactive PLD on the surface of WSe<sup>2</sup> layer.

To change the conditions under which the MoS*x*~4/WSe<sup>2</sup> interface is formed, one can employ a different technique for the deposition of a molybdenum sulfide film—one that prevents the influence of plasma-activated H2S gas. Fominski et al. [55] and Giuffredi et al. [32] demonstrate that the pulsed laser ablation of a MoS<sup>2</sup> target in a buffer gas enables the formation of MoS*<sup>x</sup>* films with an increased concentration of sulfur (*x* ≥ 3). These films have an extremely high electrocatalytic HER activity. The area of MoS*<sup>x</sup>* film deposition (i.e., WSe<sup>2</sup> nanopetals) can be oriented randomly to the axis of the plume expansion [56]. If this technique for MoS*<sup>x</sup>* film deposition is applied, the deposition of Mo and S atom flux on the interface with the WSe<sup>2</sup> film occurs almost simultaneously. This contributes to the formation of Mo-S chemical bonds in the growing film. The energy of atoms deposited during the ablation of the MoS<sup>2</sup> target in the on-axis PLD configuration is much lower than during the ablation of metallic Mo in the on-axis RPLD configuration. As a rule, the ablation of metals occurs in the conditions of effective laser plume ionization under the influence of more powerful laser pulses. This factor can also impact chemical processes at the MoS*x*/WSe<sup>2</sup> interface.

The regulation of the MoS*x*/WSe<sup>2</sup> interface formation is not the only factor that affects the efficiency of photo-assisted HER. Another one is the texture of the WSe<sup>2</sup> layer [10,13]. WSe<sup>2</sup> petals sitting along the substrate surface minimize the impact of edge states on the recombination on nonequilibrium carriers. Yet, the orthogonal orientation of the petals increases the area of the surface involved in catalysis. The negative effect of edge states can be reduced through their passivation by a MoS*<sup>x</sup>* catalyst. We carried out additional studies to obtain WSe<sup>2</sup> by off-axis PLD at varying buffer gas pressures. This factor did not have a marked effect on the texture of WSe<sup>2</sup> films. When this WSe<sup>2</sup> formation technique is used, other parameters of off-axis PLD may vary as well. These are laser fluence, the laser plume incidence angle, deposition temperature, etc. Co-deposition with some metals (for example, Pd [13]) will also affect the growth of WSe<sup>2</sup> films. The optimization of regimes for obtaining WSe<sup>2</sup> films with a targeted structure by laser-based methods is a central

condition for creating HER photocatalysts with suitable characteristics. Achieving the latter requires further research into the MoS*x*/WSe<sup>2</sup> heterojunction system.

#### **5. Conclusions**

Using different PLD configurations makes it possible to fully form a HER catalyst (in one production vessel) on a dielectric substrate (sapphire). A robust rear contact (conducting layer) was obtained using B-doped amorphous carbon by traditional onaxis PLD. A nanostructured WSe<sup>2</sup> layer was grown on the C(B) contact layer. The WSe<sup>2</sup> layer consisted of differently oriented nano-petals, which had a nearly perfect 2H-WSe<sup>2</sup> crystal lattice. To obtain a WSe<sup>2</sup> layer, off-axis PLD was performed in a buffer gas. A catalytic MoS*x*~4 layer was created on the surface of WSe<sup>2</sup> petals by on-axis reactive PLD from Mo target in H2S gas. The temperature of functional layer formation for a MoS*x*~4/WSe2/C(B)/Al2O<sup>3</sup> photocathode ranged between 22–700 ◦C.

The MoS*x*~4/WSe2/C(B)/Al2O<sup>3</sup> photocathode obtained by laser-based processes has the following characteristics as regards HER in 0.5M H2SO<sup>4</sup> acid solution during light irradiation with an intensity of 100 mW/cm<sup>2</sup> : the current density at 0 V (RHE) is ~3 A/cm<sup>2</sup> ; the onset potential reaches 400 mV (RHE). Given that these photocathodes are made from relatively cheap materials commonly found in nature, these are suitable characteristics. The performance of MoS*x*/WSe<sup>2</sup> heterojunction system for photo-assisted water splitting for hydrogen production can be substantially increased by enhancing the composition of the photocatalyst (i.e., employing a different rear contact) and optimizing PLD regimes for creating functional semiconductor layers.

**Author Contributions:** Conceptualization, V.F. and R.R.; methodology, M.D.; PLD of the films, D.F.; investigation, XPS studies, O.R.; investigation, optical properties, S.N., V.V., and N.D.; writing original draft preparation, V.F. and R.R.; writing—review and editing, V.F. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the Russian Science Foundation, grant number 19-19-00081.

**Data Availability Statement:** Not applicable.

**Acknowledgments:** Sample characterization by optical spectroscopy has been done in Center for Photonics and 2D Materials, MIPT. V.V., S.N., and N.D. gratefully acknowledge financial support from the Ministry of Science and Higher Education of the Russian Federation (No. 0714-2020-0002).

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


## *Article* **Synthesis of Air-Stable Cu Nanoparticles Using Laser Reduction in Liquid**

**Ashish Nag , Laysa Mariela Frias Batista and Katharine Moore Tibbetts \***

Department of Chemistry, Virginia Commonwealth University, Richmond, VA 23284, USA; nagab@vcu.edu (A.N.); friasbatistlm@vcu.edu (L.M.F.B.)

**\*** Correspondence: kmtibbetts@vcu.edu

**Abstract:** We report the synthesis of air-stable Cu nanoparticles (NPs) using the bottom-up laser reduction in liquid method. Precursor solutions of copper acetlyacetonate in a mixture of methanol and isopropyl alcohol were irradiated with femtosecond laser pulses to produce Cu NPs. The Cu NPs were left at ambient conditions and analyzed at different ages up to seven days. TEM analysis indicates a broad size distribution of spherical NPs surrounded by a carbon matrix, with the majority of the NPs less than 10 nm and small numbers of large particles up to ∼100 nm in diameter. XRD collected over seven days confirmed the presence of fcc-Cu NPs, with some amorphous Cu2O, indicating the stability of the zero-valent Cu phase. Raman, FTIR, and XPS data for oxygen and carbon regions put together indicated the presence of a graphite oxide-like carbon matrix with oxygen functional groups that developed within the first 24 h after synthesis. The Cu NPs were highly active towards the model catalytic reaction of *para*-nitrophenol reduction in the presence of NaBH<sup>4</sup> .

**Keywords:** laser synthesis; laser reduction in liquid; copper nanoparticles; para-nitrophenol

**Citation:** Nag, A.; Frias Batista, L.M.; Tibbetts, K.M. Synthesis of Air-Stable Cu Nanoparticles Using Femtosecond Laser Reduction in Liquid. *Nanomaterials* **2021**, *11*, 814. https:// doi.org/10.3390/nano11030814

Academic Editor: Mohamed Boutinguiza

Received: 15 March 2021 Accepted: 21 March 2021 Published: 23 March 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

#### **1. Introduction**

Laser synthesis techniques have emerged over the last decade as reliable methods for producing pure nanomaterials (NMs) [1–3]. Laser synthesis has several advantages over traditional chemical synthesis methods including: facile generation of metastable phases and bonding environments, rapid conversion of precursors to NM products, and avoidance of capping ligands. For instance, alloys of immiscible metals can be formed [4], reactions can be completed in seconds [5], and grams per hour synthesis yields can be attained with high-repetition rate lasers [6]. Moreover, laser synthesis techniques are considered 'green' because they do not employ toxic chemical reducing agents or surfactants and produce little chemical waste [7].

The lack of otherwise required capping ligands results in high purity nanoparticles that are important to several applications. Biofunctional assemblies can be produced by functionalizing laser-generated nanoparticles (NPs) for in-vitro applications [8]. The absence of ligands makes more active sites available for catalysis, resulting in laser-synthesized NPs often having higher activities than their conventional counterparts [9–11]. Higher purity also makes laser-synthesized NPs attractive candidates for other biomedical applications and as references for modeling chemical reactions [12–15]. In addition, photoluminescence can be introduced to the NPs by in-situ generation of carbon shells through laser-induced decomposition of organic solvents [16,17].

Cu NPs in particular are of high interest both due to the natural abundance of copper and their myriad applications in catalysis, electronics, and biology [18–22]. However, conventional wet-chemical methods used to synthesize Cu-based NPs require toxic solvents, reducing agents, or both [23–25]. Hence, greener synthesis routes to Cu NPs are of primary importance. To address this need, pulsed laser ablation in liquid (PLAL) has been widely employed to generate Cu NPs [26–30]. In these syntheses, copper oxides are a major product when ablation is conducted in water [26–29], and Cu<sup>0</sup> phases are only

stable when the ablation liquid contains an organic solvent, such as acetone, methanol, or ethanol [28–30].

The *top-down* PLAL method is by far the most common laser synthesis technique used to produce colloidal metal NPs [1]. PLAL involves focusing of laser beam on the surface of a solid or powdered target immersed in a solvent, causing the removal of target material from the surface and its coalescence into colloidal NPs. However, NP products from PLAL in many cases exhibit bimodal size distributions due to the ejection of both small clusters and large droplets from the surface [31–33]. An alternative method is *bottom-up* laser reduction in liquid (LRL), which involves focusing picosecond (ps, 10−<sup>12</sup> s) or femtosecond (fs, 10−<sup>15</sup> s) laser pulses into a solution of molecular precursors to generate a dense plasma containing electrons that reduce metal ions to colloidal NPs. LRL can enable superior control over Au NP sizes in a single step when the chemistry of the precursor solution and laser irradiation conditions are carefully controlled [34,35].

A current limitation to the wide use of LRL is that the vast majority of studies reported to date focus on the easily reduced noble metals Au [34–38], Ag [39–41], Pt [42], and their alloys [43–45] because these metals are resistant to oxidation. Unlike in PLAL where the zero-valent metal is present in the initial target, metal NP formation in LRL requires chemical reactions between the molecular precursor(s) and the reactive species in LRL plasma. In aqueous solution, the major species are hydrated electrons (e – aq ) and hydroxyl radicals (OH · ) [38]. Whereas e – aq are exceptionally strong reducing agents towards metal ions, OH · radicals and their recombination product H2O<sup>2</sup> can back-oxidize zero-valent metal atoms. Although Au is resistant to back-oxidation and HAuCl<sup>4</sup> precursor can be reduced by H2O<sup>2</sup> [37], effective production of Ag-containing NPs by LRL requires the addition of OH · scavengers, such as ammonia or isopropyl alcohol, to prevent backoxidation of Ag [40,41,45]. Moreover, the few LRL studies on the non-noble metal Fe report production of only Fe oxides [46,47].

The detrimental effects of reactive oxygen species on Cu NP synthesis have been widely known since Dhas et al. [48] reported that OH · and H2O<sup>2</sup> formed during sonochemical synthesis lead to the formation of copper oxides. In PLAL of Cu metal in water, Cu oxides are formed due to the presence of OH · radicals in the laser plasma from water, dissolved O<sup>2</sup> , or both [28]. Oxide formation during PLAL can be hindered through the use of organic solvents [27–30], which also can result in the formation of a protective carbon shell around the Cu NPs [27,28]. On the basis of these results, we designed an air-free LRL synthesis route to Cu NPs using an isopropyl alcohol/methanol solvent mixture. We report, to the best of our knowledge, the first LRL synthesis of air-stable Cu NPs and their catalytic activity towards reduction of *para*-nitrophenol in the presence of NaBH<sup>4</sup> .

#### **2. Materials and Methods**

#### *2.1. Materials*

Copper acetylacetonate (Cu(acac)<sup>2</sup> , Acros Organics, Fair Lawn, NJ, USA), isopropyl alcohol (IPA, Fisher Scientific, Waltham, MA, USA), methanol (MeOH, Fisher Scientific, Waltham, MA, USA), sodium borohydride (NaBH<sup>4</sup> , Acros Organics, Fair Lawn, NJ, USA), and *para*-nitrophenol (PNP, Acros Organics, Fair Lawn, NJ, USA) were used as received.

#### *2.2. Cu NP synthesis*

A working solution of 0.8 mM Cu(acac)<sup>2</sup> in 25%/75% (*v/v*) MeOH/IPA solvent was purged with nitrogen for 20 min before 3.0 mL of the solution was transferred to a nitrogenfilled 10 × 10 × 40 mm quartz fluorimeter cuvette. The solution was then irradiated with laser pulses to yield Cu NPs. Upon formation, Cu NPs were exposed to air and analyzed at different time intervals (0 h, 3 h, 24 h, and 7 days) to study the aging of NPs. By 24 h after synthesis, all the Cu NPs are precipitated out of the solution but can be redispersed via brief sonication.

#### *2.3. Instrumentation*

#### 2.3.1. Laser Synthesis

The experimental setup is described elsewhere [38]. Briefly, samples were irradiated using a commercial titanium-sapphire chirped-pulse amplifier (Astrella, Coherent, Inc., Santa Clara, CA, USA), delivering 30 fs pulses, with the bandwidth centered at 800 nm and a repetition rate of 1 kHz; 2 mJ laser pulses were focused using a *f* = 5 cm aspheric lens into the center of the sample cuvette to produce a peak irradiance of <sup>5</sup> <sup>×</sup> <sup>10</sup><sup>16</sup> Wcm−<sup>2</sup> . Details for calculating peak irradiance can be found in Ref. [38]. The sealed cuvettes containing Cu(acac)<sup>2</sup> solution were irradiated for 10 min.

#### 2.3.2. UV-Vis Spectroscopy

Conversion of precursors to Cu NPs was monitored using a home-built in-situ UV-vis spectrophotometer described in Ref. [38] and the irradiation was stopped when no further growth in Cu surface plasmon resonance (SPR) was observed. The catalytic performance of Cu NPs was tested for the reduction of PNP with NaBH<sup>4</sup> (see Section 2.4) by employing a second home-built in-situ UV-vis spectrophotometer described in Ref. [49]. Finally, to observe the effect of aging of Cu NPs, absorbance data was recorded using Agilent 8453 UV-vis spectrophotometer at times up to 7 days after synthesis.

#### 2.3.3. Transmission Electron Microscopy (TEM)

Cu NPs were visualized using TEM (JEOL JEM-1230 TEM) at 120 kV. A diluted solution of colloidal Cu NPs was drop-casted onto carbon-coated grids (Structure Probe, Inc., West Chester, PA, USA) and left to dry for 24 h or longer. Average sizes and size distributions were measured using ImageJ software. At least 350 particles from images of three separate areas of a TEM grid were evaluated.

#### 2.3.4. X-ray Photoelectron Spectroscopy (XPS)

XP spectra were collected on a PHI VersaProbe III Scanning XPS Microprobe with a monochromated Al K*α* X-ray source (1486.6 eV), with a typical resolution of 0.4–0.5 eV. Survey scans and high resolution scans were collected with pass energies of 280 eV and 26 eV, respectively. Charge neutralization was done by running an ion gun and a flood gun during sample analysis. The measurement spot diameter was 200 µm with take off angle of 90◦ and the detector at 45◦ . Spectral analysis was carried out using PHI Multipak XPS software with 70% Gaussian/Lorentzian convolution to fit each spectral peak. Samples were prepared by drop-casting Cu NPs on a gold-sputtered silicon wafer, followed by drying under vacuum for at least 24 h. All spectra were corrected using a Au4f peak shift to center at 84.0 eV.

#### 2.3.5. FTIR Spectroscopy

A Thermo Scientific Nicolet iS50 FTIR spectrometer equipped with a mid- and far-IRcapable diamond ATR was used to record FTIR spectra. All spectra were obtained using 32 scans in the range from 4000 to 400 cm−<sup>1</sup> with 4 cm−<sup>1</sup> resolution. Cu NPs were directly drop-casted on to the diamond crystal.

#### 2.3.6. X-ray Diffraction (XRD)

XRD data was collected on a Panalytical Empyrean Diffractometer with CuK*α* radiation (0.15418 nm) at 40 kV and 45 mA, with scanning angle (2*θ*) of 30–80◦ and a gonio focusing geometry. Sample preparation involved drop-casting of Cu NPs on a low-background silicon substrate, followed by drying under vacuum for at least 24 h.

#### 2.3.7. Raman Spectroscopy

Raman spectra were recorded on a Thermo Scientific DXR3 SmartRaman Spectrometer involving a 532 nm excitation laser. All spectra obtained were averaged over 32 scans in the range of 50 to 3500 cm−<sup>1</sup> with 5 cm−<sup>1</sup> resolution. Samples were prepared by drop-casting Cu NPs on a silicon wafer, followed by drying under vacuum for at least 24 h.

#### *2.4. Catalytic Reduction of Para-Nitrophenol (PNP)*

PNP reduction reactions were carried out in a home-built in situ UV-vis spectrometer setup described elsewhere [35] at different time intervals after laser synthesis (0 h, 3 h, 24 h, and 7 days). Spectra were recorded every 2.4 s using LabVIEW software (National Instruments). A solution containing a final concentration of 0.1 mM PNP and 10 mM of freshly prepared NaBH<sup>4</sup> was prepared in a 10 × 10 × 40 mm quartz fluorimeter cuvette with a magnetic stir bar, resulting in the formation of *p*-nitrophenolate ion with UV-vis absorbance at 400 nm. Prior to the addition of the catalyst, this peak was observed for 24 s to confirm no reaction occurred in the absence of a catalyst. After this period, 300 µL of the Cu NP catalyst was added, triggering the reduction of PNP. Data collection was terminated when the absorbance at 400 nm (*p*-nitrophenolate ion) had disappeared.

#### **3. Results**

#### *3.1. Physical Characterization*

Figure 1 shows TEM images and size distributions for Cu NPs samples obtained 0 h (a), 3 h (b), and 24 h (c) after laser synthesis. In all images, small <10 nm NPs surrounding significantly larger NPs are observed. Fitting the size distributions to a log-normal function primarily captured the small NPs, producing mean diameters of 3.30 nm (PDI = 0.16), 4.54 nm (PDI = 0.19), and 4.57 nm (PDI = 0.16) for 0 h, 3 h, and 24 h samples, respectively. Despite the large numbers of small NPs, they make up a tiny fraction of the overall mass of Cu NPs, as evident in the mass-weighted size distributions. Nevertheless, the contribution of NPs smaller than 10 nm to the overall mass slightly increased from 0.16 % at 0 h to 0.34 % at 3 h and 24 h. The slight size increase in the small NPs after the initial synthesis was also observed for Fe oxide NPs synthesized by LRL from Fe(acac)<sup>3</sup> in hexane [47], although, in that work, no large NPs were observed. A carbon shell around clusters of Cu NPs is also visible in Figure 1, particularly for the 3 h and 24 h samples. Similar carbon shells have been observed for Cu NPs produced by PLAL in acetone [28] and methanol [27].

**Figure 1.** TEM images and corresponding number-weighted and mass-weighted size distributions for samples analyzed (**a**) 0 h, (**b**) 3 h, and (**c**) 24 h after synthesis. Scale bars in magnified regions are 20 nm.

Figure 2 shows the XRD spectrum of the 24 h sample. Intense diffraction peaks observed at 43.35◦ , 50.47◦ , and 74.18◦ were indexed to (111), (200), and (220) planes of fcc-Cu (JCPDS no. 01-085-1326), respectively. A broad feature centered at 36.85◦ indexed to cubic-Cu2O(111) (JCPDS no. 00-005-0667) was also observed. Using the Scherrer equation, the crystallite sizes were calculated to be 29.89 nm for the Cu(111) [50]. Thus, we can infer that the large NPs observed in Figure 1 are Cu metal NPs, whereas Cu2O exists as an amorphous thin layer around the Cu NPs, as the smaller NPs, or both.

**Figure 2.** XRD spectrum of 24 h sample (blue), with references for Cu2O (cyan) and Cu (dark red). Inset magnifies a broad peak observed in the spectra for 24 h sample indexed to Cu2O (111).

Figure 3 shows XP spectra obtained for the 24 h sample. Two peaks in the Cu2p3/2 region (Figure 3a) at 933.0 eV (cyan) and 934.2 eV (violet) were assigned to either Cu<sup>0</sup> or Cu<sup>+</sup> and Cu2+, respectively [51]. To distinguish between Cu<sup>0</sup> and Cu+, the CuLMM region was analyzed (Figure 3b) and resulted in two peaks at 568.4 eV (blue) and 570.6 eV (magenta) assigned to Cu<sup>0</sup> and Cu+, respectively [51]. The presence of both Cu<sup>0</sup> and Cu<sup>+</sup> in the XP spectra is consistent with the XRD results (Figure 2). In the O1s region (Figure 3c), the peaks at 530.8 eV (magenta), 531.7 eV (light blue), 532.2 eV (light green), and 533.0 eV (dark green) were assigned to Cu–O, C=O, –OH and C–O species, respectively [52,53]. In the C1s region (Figure 3d), peaks at 284.4 eV (red), 284.8 eV (orange), 286.1 eV (dark green), and 288.4 eV (light blue) are assigned to C=C, C–C, C–O, and C=O species, respectively [53]. This collection of carbon species is associated with the carbon shell observed around the Cu NPs (Figure 1).

**Figure 3.** XPS spectra of (**a**) Cu2p3/2, (**b**) CuLMM, (**c**) O1s and (**d**) C1s regions with fitted peaks for 24 h samples.

Spectral characterization of the Cu NPs (Figure 4) confirms the chemical species assigned in Figures 2 and 3. Figure 4a shows the UV-vis absorbance data collected at 0 h, 3 h, and 24 h after synthesis. An absorbance peak at 574 nm assigned to the Cu surface plasmon resonance (SPR) was observed at 0 h, followed by broadening and red-shifting to 586 nm and 596 nm at 3 h and 24 h, respectively. These changes in the absorbance peaks are associated with the oxidation of Cu to Cu2O over time [54] and likely indicate the formation

of a Cu2O shell around the large Cu NPs on the basis of the XRD spectrum (Figure 2). Raman spectra shown in Figure 4b closely resemble graphite oxide spectra [55]. The intense G band at 1600 cm−<sup>1</sup> and D band at 1350 cm−<sup>1</sup> correspond to the in-plane vibrational modes of sp<sup>2</sup> -hybridized carbon atoms and structural disorder due to functionalization, respectively [56]. Moreover, the weak and broad 2D and D + G bands at 2680 cm−<sup>1</sup> and 2910 cm−<sup>1</sup> arise from disorder due to formation of oxygen functional groups [55]. The presence of oxygen functional groups is confirmed by the FTIR spectra (Figure 4c) with peaks indicating C=O stretch of COOH groups at 1722 and 1696 cm−<sup>1</sup> , COO− stretch at 1415, 1522, 1575 cm−<sup>1</sup> , O–H deformations of C–OH groups at 1356 cm−<sup>1</sup> , C–O stretch of epoxide groups at 1261 cm−<sup>1</sup> , C–O stretch at 1000–1100 cm−<sup>1</sup> , and C–H bend at 800 cm−<sup>1</sup> [56,57], although some of these bands may be overlapped with vibrational modes of other functional groups. Notably, these IR peaks grow in intensity over the course of 24 h after synthesis, suggesting that the carbon shell formation occurs mostly after laser irradiation is terminated. Overall, these data and the XP spectra for C1s and O1s (Figure 3a,b) indicate that the carbon shell around the Cu NPs consists of disordered graphite oxide-like structures that contain multiple different functional groups.

**Figure 4.** (**a**) UV-vis absorbance spectra for 0 h, 3 h, 24 h samples and precursor. (**b**) Raman spectra for 24 h samples. (**c**) FTIR spectra for 0 h, 3 h, 24 h samples and precursor.

#### *3.2. Catalytic Activity of Cu NPs*

Reduction of PNP by NaBH<sup>4</sup> is a commonly used model reaction to test catalytic activity of metal NPs by employing UV-vis spectroscopy to monitor the decrease in absorbance of the *p*-nitrophenolate ion at 400 nm, which allows for convenient determination of pseudo-first-order rate constants when excess NaBH<sup>4</sup> is present [58,59]. Figure 5a shows representative kinetic curves obtained from the ratio of the natural log of the 400 nm absorbance feature at time *t*, A(*t*), to the initial absorbance, A(0), as a function of reaction time *t* for the 0 h, 3 h, and 24 h samples. The slopes obtained from the linear regions of the

curves (shown in black) give the apparent rate constants (*k*app). The values of *k*app were averaged over three different samples for all three post-synthesis times and then converted to the mass-specific rate constants (*k*) of 1084, 1479, and 1927 s−1g −1 for 0 h, 3 h, and 24 h samples (Figure 5b). The specific rate at 3 h is comparable to the rate of 1490 s−1g −1 for Cu NPs synthesized by PLAL in a water-ethanol mixture [30]. Although the rate constant is higher at 24 h, this increase could be due to some evaporation of the solvent over time, resulting in concentration of the Cu NPs and an artificially high measured rate.

**Figure 5.** (**a**) Time dependence of the absorbance of *p*-nitrophenolate ions at 400 nm for a representative sample at 0 h, 3 h, and 24 h; black portion of line is where k*app* (s−<sup>1</sup> ) is extracted. (**b**) Specific rate constants (s−1g −1 ) with error bars generated over three runs for 0 h, 3 h, 24 h samples each.

#### *3.3. Stability of Cu NPs*

Figure 6 compares the UV-vis absorbance, XRD spectrum, and PNP kinetics for 1 day (24 h) and 7 day samples. The Cu SPR peak observed for 7 day sample was only slightly red shifted compared to 1 day sample (Figure 6a), indicating no significant oxidation of Cu NPs. No changes were observed in the XRD peaks for 7 day samples (Figure 6b) compared 1 day samples (Figure 2), further confirming no or insignificant oxidation of Cu NPs. Accordingly, similar rate constants were obtained for 1 and 7 day samples for the PNP reduction reaction (Figure 6c). Collectively, these data indicate that the LRL Cu NPs are highly stable to oxidation.

**Figure 6.** (**a**) UV-vis absorbance spectra measured at 1 day (blue) and 7 days (magenta). (**b**) XRD spectra for 7 day sample. (**c**) Time dependence of the absorbance of *p*-nitrophenolate ions at 400 nm for 1 day (blue) and 7 days (bright pink) samples; black portion of line is where *k*app (s−<sup>1</sup> ) is extracted.

#### **4. Discussion**

The production of Cu metal NPs using LRL requires both an organic solvent and air-free conditions, as no laser-induced conversion of Cu(acac)<sup>2</sup> precursor was observed when the solvent was water or when the precursor in IPA/methanol mixture was irradiated under ambient conditions. This result underscores the importance of minimizing reactive oxygen species formation during LRL, which was also needed to effectively control the sizes of Au NPs [35] and enable formation of Ag-containing NPs [41,44,45]. The important role of reactive oxygen species has also been noted in PLAL synthesis of Cu NPs, where air removal completely eliminated formation of the highly oxidized CuO phase during PLAL in water [28]. Collectively, these results highlight similarities in the chemical reaction pathways induced during both PLAL and LRL that can lead to metal NP oxidation, as well as the potential for the same strategies to mitigate oxide formation.

The formation of substantial quantities of sub-10 nm Cu NPs in our LRL synthesis is similar to previous LRL results for Au and Ag NPs using well-controlled solvent chemistry [35,45], whereas the large Cu NPs up to ∼100 nm in diameter resemble the large Fe oxide NPs obtained from LRL of ferrocene in hexane [46]. The formation of two distinct size distributions of Cu NPs strongly suggests the participation of two different reaction mechanisms [33], although specific identification of these mechanisms is beyond the scope of this work. Nevertheless, on the basis of the PLAL literature demonstrating that substantial control over Cu NP size and morphology can be achieved by changing solvent mixtures [27–30], we anticipate that further exploration of different solvents will enable better control over Cu NP sizes using LRL. Finally, we note that the catalytic activity of the LRL Cu NPs to PNP reduction is comparable to that reported for PLAL Cu NPs [30], despite the presence of a thick carbon shell around our NPs. This result indicates that the carbon shell is sufficiently permeable to allow for catalytic reactions to take place and suggests that our LRL Cu NPs may have additional catalytic applications in areas, such as CO<sup>2</sup> reduction or cross-coupling.

The formation of a substantial carbon shell around the LRL Cu NPs after termination of laser irradiation strongly suggests that the carbon formation is not entirely attributable to direct laser-induced solvent decomposition, as with PLAL in organic solvents [16,17,27,28]. Moreover, the lack of carbon shell formation during LRL of Fe(acac)<sup>3</sup> in water [47] indicates that the acetylacetonate ligands from the precursor are not sufficient to induce carbon shell formation in LRL. On the basis of these results, we can speculate that the carbon shell is formed from catalytic activation on the Cu NP metal surfaces of long-lived solvent and ligand byproducts produced during laser irradiation. Cu metal is known to be a highly active catalyst in carbon-carbon cross-coupling reactions [19] and induces graphene growth from aliphatic alcohols under low-temperature CVD conditions [60]. Hence, we anticipate that the active bare Cu surfaces present immediately following laser irradiation catalyze the formation of the carbon shell. Although carbon shells are widely known to effectively protect transition metal NPs from oxidation [61–63], the evident permeability of the carbon around the LRL Cu NPs on the basis of their high catalytic activity suggests that a thin Cu2O layer around the Cu NPs also contributes to their observed stability [61].

#### **5. Conclusions**

We have reported the first synthesis of air-stable Cu NPs using the laser reduction in liquid (LRL) approach. Both small (<10 nm) and large (up to 100 nm) spherical NPs were observed, which primarily consisted of Cu<sup>0</sup> metal on the basis of XRD and XPS analysis, although a Cu2O shell around the large NPs may be present. The Cu NPs exhibit remarkable stability over 7 days on the basis of the lack of significant changes observed in the UV-vis absorbance and XRD features and the similar rate constants obtained for PNP reduction. The LRL Cu NPs compare favorably with PLAL-synthesized Cu NPs in terms of Cu<sup>0</sup> content, stability, and catalytic activity. The insights into LRL of Cu2<sup>+</sup> gained in this manuscript can be extended to other metals for which oxidation of NPs is commonly observed during PLAL or LRL. Finally, further development of the synthesis conditions, such as conducting LRL in a flow setup, could increase Cu NPs yield to produce heterogeneous catalysts for applications, such as cross-coupling reactions and CO<sup>2</sup> reduction.

**Author Contributions:** K.M.T. and A.N. designed the research. A.N. performed the experiments and analyzed the data. L.M.F.B. performed experiments and processed the TEM data. A.N. and K.M.T. wrote the manuscript. All authors have read and agreed to the published version of the manuscript.

**Funding:** This work was supported by the American Chemical Society Petroleum Research Fund through Grant 57799-DNI10. L.M.F.B. acknowledges generous financial support from an Altria Graduate Research Fellowship.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** Data presented in this study are available on request from the corresponding author.

**Acknowledgments:** Microscopy was performed at the VCU Department of Anatomy and Neurobiology Microscopy Facility, supported by the Higher Education Equipment Trust Fund Grant No. 236160307. We would like to acknowledge the VCU Nanomaterials Core Characterization Facility and Chemistry Instrumentation Facility for additional characterization.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

