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Chalcogenide perovskite $\mathtt{B a Z r S}_{3}$ thin-film electronic and optoelectronic devices by low temperature processing

Zhonghai $\mathrm{Yu}^{,\mathrm{a,1}}$ , Xiucheng Wei b,1, Yixiong Zheng c,1, Haolei Hui b, Mengying Bian b, Samyak Dhole c, Jung-Hun Seo c,*, Yi-Yang Sun d, Quanxi Jia c, Shengbai Zhang e, Sen Yang a,*, Hao Zeng b,*

a MOE Key Laboratory for Nonequilibrium Synthesis and Modulation of Condensed Matter, Xi’an Jiaotong University, Xi’an 710049, China
b Department of Physics, University at Buffalo, The State University of New York, Buffalo, NY 14260, USA
c Department of Materials Design and Innovation, University at Buffalo, The State University of New York, Buffalo, NY 14260, USA
d State Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 201899,
China

e Department of Physics, Applied Physics & Astronomy, Rensselaer Polytechnic Institute, Troy, NY 12180, USA

A R T I C L E I N F O

Keywords:
Chalcogenide perovskite
Pulsed laser deposition
Photodetector
Field effect transistor

A B S T R A C T

Owing to its superior visible light absorption and high chemical stability, chalcogenide perovskite barium zir­ conium sulfide $\mathrm{(BaZrS_{3})}$ has attracted significant attention in the past few years as a potential alternative to hybrid halide perovskites for optoelectronics. However, the high processing temperatures of $\mathrm{BaZrS}{3}$ thin films at above $1000,^{\circ}\mathrm{C}$ severely limits their potential for device applications. Herein, we report the synthesis of $\mathrm{\bfBaZrS{3}}$ thin films at temperatures as low as $500,^{\circ}\mathrm{C}$ , by changing the chemical reaction pathway. The single phase $\mathrm{\bfBaZrS}_{3}$ thin film was confirmed by X-ray diffraction and Raman spectroscopy. Atomic force microscopy and scanning electron microscopy show that crystalline size and surface roughness were consistently reduced with decreasing annealing temperature. The lower temperatures further eliminate sulfur vacancies and carbon contaminations associated with high temperature processing. The ability to synthesize chalcogenide perovskite thin films at lower temperatures removes a major hurdle for their device fabrication. The photodetectors demonstrate fast response and an on/off ratio of 80. The fabricated field effect transistors show an ambipolar behavior with electron and hole mobilities of $16.8,\mathrm{cm}^{2}/\mathrm{Vs}$ and $2.6,\mathrm{cm}^{2}/\mathrm{Vs},$ , respectively.

1. Introduction

Organic-inorganic hybrid halide perovskites have taken the center stage in photovoltaics research over the past decade [1–4], since Kojima et al. first introduced their application as a visible-light sensitizer in 2009 [5]. Many studies have further established their outstanding per­ formance for optoelectronic applications, such as photodetectors [6,7], light emitting diodes (LEDs) [8,9], optical sensors [10,11], and espe­ cially solar cells [12,13]. However, the best performing halide perov­ skites for solar cells often contain the toxic element Pb [5,14]. Furthermore, hybrid perovskites show inherent instability against moisture, heat, light illumination, and electric field. These disadvan­ tages have severely impeded their commercial applications. A lot of effort has been devoted to enhancing the stability of organic perovskites, such as A-site cation engineering (e.g. replacing $\mathbf{M}\mathbf{A}^{+}$ by $\mathbf{F}\mathbf{A}^{+}$ and $\mathsf{C}_{S}^{+}.$ ) [15–17] and dimension reduction (e.g. using the two-dimensional Ruddlensden-Popper phase) [18–20].

To address the toxicity and stability limitations of hybrid perov­ skites, inorganic chalcogenide perovskites have been proposed as an alternative family of perovskite semiconductors for optoelectronics. In 2015, Sun et al. theoretically screened $18\ \mathrm{ABX}{3}$ chalcogenide perov­ skites for photovoltaics, with A as $\mathrm{Ca/Sr/Ba}$ , B as $\mathrm{Ti/Zr/Hf{\mathrm{:}}}$ and X as S/Se [21]. Several $\mathrm{ABX}{3}$ chalcogenide perovskites were identified to have suitable band gaps and absorption properties for photovoltaics. As a prototypical chalcogenide perovskite, $\mathtt{B a Z r S}{3}$ possesses a direct band gap of $1.8,\mathrm{eV}$ , and strong near edge absorption [22–26]. Furthermore, $\mathrm{BaZrS}{3}$ has been shown to possess a stable perovskite structure against high pressure, moisture, and heat [23,27]. However, little is known about its carrier transport properties as the majority of publications are on bulk powder or single crystals. We first reported the synthesis of $\mathrm{BaZrS}{3}$ thin films by high temperature $({\sim}1000,^{\circ}\mathrm{C})$ sulfurization of $\mathrm{BaZrO}{3}$ films deposited via pulsed laser deposition (PLD) [28]. These films show $n$ -type conductivity due to sulfur vacancies, with a carrier mobility of around $10,\mathrm{cm}^{2}/\mathrm{Vs}$ . Comparotto et al. fabricated $\mathrm{BaZrS}{3}$ thin films by co-sputtering of BaS and $Z\mathbf{r}$ targets, followed by a thermal treatment at temperatures varying from 650 to $1000,^{\circ}!\mathrm{C}$ . It was reported that $900,^{\circ}\mathrm{C}$ was required for good crystallization. However, no transport properties were reported [29]. Gupta et al. synthesized $\mathrm{BaZrS}_{3}$ thin films by sulfurization of $\mathrm{{BaZrO3}}$ films made by chemical solution deposition at $1050,^{\circ}\mathrm{C}$ [30]. Such high processing temperatures are incompatible with device fabrication, and could also lead to unintended high concentration of electron doping due to sulfur vacancy defects. Thus, lowering the processing temperature of chalcogenide perovskite thin films is a crucial step for establishing this family of materials as a viable competitor to hybrid organic perovskites for practical applications.

In this paper we report the synthesis of $\mathrm{BaZrS}{3}$ thin films at tem­ peratures as low as $500,^{\circ}\mathrm{C}$ , which is comparable to the processing temperature of copper indium gallium selenide (CIGS) solar cells [31, 32]. This is achieved by changing the chemical reaction pathway from sulfurization of oxide perovskites to crystallization of pulsed laser deposited amorphous $\mathrm{\ttBaZrS{x}}$ films. The films processed at low temper­ atures show stoichiometric single-phase orthorhombic perovskite structure with smooth surfaces. Elimination of excess sulfur vacancies leads to reduced carrier concentration to $\sim!4\times10^{10}\ensuremath{~\mathrm{\mm}^{-3}}$ . As a demonstration for the potential use of the $\mathrm{BaZrS}{3}$ thin films towards various electronic and optoelectronic applications, $\mathrm{BaZrS}{3}$ thin-film photodetectors and field effect transistors (FET) were fabricated. The photodetector measurements show photocurrent close to 2 orders of magnitude greater than dark current for the films annealed at $650,^{\circ}\mathrm{C}$ . The FETs with ionic liquid (IL) gating show ambipolar behavior, with an on/off ratio of $6.5\times10^{3}$ at gate voltage of $_{1,\mathrm{V}}$ and carrier mobilities estimated to be $16.8,\mathrm{cm}^{2}/\mathrm{Vs}$ and $2.6,\mathrm{cm}^{2}/\mathrm{Vs}$ for electrons and holes, respectively. Our study paves the way for the integration of chalco­ genide perovskite thin films in optoelectronic devices.

2. Experimental

2.1. Preparation of the BaZrS3 PLD target

$\mathrm{BaZrS}{3}$ powder was synthesized from $\mathtt{B a Z r O3}$ powder following published procedure [24]. The pre-synthesized $\mathtt{B a Z r S}{3}$ powder was cold-pressed into a pellet with a $25,\mathrm{T}$ hydraulic press using a circular mold of $20,\mathrm{mm}$ in diameter. Before pressing, a small amount of polyvinyl alcohol (PVA) was added to enhance the adhesion and stability of the pellet. The $\mathrm{BaZrS}_{3}$ pellet was sintered at $1050,^{\circ}\mathrm{C}$ in Ar atmosphere for $2,\mathrm{{h}}$ , and used as target for PLD.

2.2. Synthesis of $B a Z r S_{3}$ thin films

A $1,\mathrm{cm}\times1,\mathrm{cm}\ (0001)$ sapphire wafer was loaded into the ultrahigh vacuum (UHV) deposition chamber of the PLD system equipped with a KrF pulsed excimer laser $(\lambda,{=},248,\mathrm{nm})$ , at a base pressure of $1\times{10}^{-8}$ torr. The substrate holder was kept at temperatures of $500{-}700,^{\circ}!\mathsf{C}$ during the deposition, with a laser repetition rate of $5,\mathrm{Hz}$ and laser fluence of $5,\mathrm{J/cm^{2}}$ . During the deposition, the substrate and target were both rotated at $30,\mathrm{rpm}$ to ensure homogeneity of the films. The as-grown amorphous $\mathtt{B a Z r S_{X}}$ thin films were cooled down to room temperature in the UHV chamber and then transferred to an MTI two-zone tube furnace in a 2-inch diameter quartz tube for annealing under $\mathrm{C}S_{2}$ carried by Ar gas flow at 12 standard cubic centimeters (sccm). The heating profile and a sketch of the setup are shown in the Supporting Information (SI) Fig S1. The as-deposited thin films were placed at the center of zone 2 downstream, with temperatures ranging from 500 to $900,^{\circ}\mathrm{C}$ and corre­ sponding total annealing time ranging from 2 to $^{8,\mathrm{h}}$ . The temperature of zone 1 was kept at $900,^{\circ}\mathrm{C}$ for complete decomposition of $\mathrm{C}!S_{2}$ as the decomposition rate of $C S_{2}$ is highly temperature dependent [33].

2.3. Thin film characterizations

The X-ray Diffraction (XRD) θ-2θ scans were performed using an X′pert Pro X-ray diffractometer operating at $1.6,\mathrm{kW}$ (Cu Kα radiation). Raman spectra were obtained from a HORIBA Raman spectrometer working under $532,\mathrm{nm}$ laser excitation. The scanning electron micro­ scopy (SEM) images and energy dispersive X-ray spectroscopy (EDX) analysis were acquired from a Focused Ion Beam-Scanning Electron Microscope (FIB-SEM) – Carl Zeiss AURIGA CrossBeam with an Oxford EDS system. The Bruker AXS-Innova atomic force microscope (AFM) was used to obtain the surface roughness and thickness profiles. The ab­ sorption spectra were collected from a Cary series UV-Vis-NIR spectro­ photometer measured from $400,\mathrm{nm}$ to $800,\mathrm{nm}$ .

2.4. Device fabrication and characterization

The photodetectors were fabricated by depositing the source and drain electrodes by electron beam evaporation of a $30,\mathrm{nm}$ thick Au film through a shadow mask, with a deposition rate of $0.4,\mathring\mathrm{A/s}$ , using an AJA Dual ATC-Orion UHV system with a base pressure of $1\times{10}^{-8}$ torr. The photodetector has a channel width of $1000,\upmu\mathrm{m}$ and a gap of $50,\upmu\mathrm{m}$ . The I$\mathtt{V}$ curves were measured by a Keithley 4200-SCS Semiconductor Parameter Analyzer, in the dark and under illumination with an exci­ tation wavelength of $532,\mathrm{nm}$ and light intensity of $0.14,\mathrm{W/cm}^{2}$ The FET devices were fabricated by the same technique through a T-shaped mask. The FET has a channel length of $50,\upmu\mathrm{m}$ and a width of $1000,\upmu\mathrm{m}$ . The side gate is separated from the source and drain electrodes by a $150,\upmu\mathrm{m}$ gap. A drop of the IL Diethylmethyl(2-methoxyethyl) ammoniumbis(trifluoromethylsulfonyl)imide (DEME-TFSI), about $20,\upmu\mathrm{L}$ , was dispensed to cover the active area of the $\mathrm{BaZrS}{3}$ FET. The device was baked at $120,^{\circ}\mathrm{C}$ for $12,\mathrm{{h}}$ in a $\mathtt{N}{2}$ glove box to remove any moisture in the IL. The drain-source current $\mathrm{I_{DS}}$ as a function of drainsource voltage $\mathrm{V_{DS}}$ under different gate voltages $\mathrm{V}{\mathrm{G}}$ and drain-gate capacitance CDG vs drain-gate voltage $\mathrm{V{G}}$ were measured by a Keithley 4200-SCS Semiconductor Parameter Analyzer.

3. Results and discussions

Conventional semiconductors and oxide/halide perovskites repre­ sent two ends of the spectrum in terms of the covalency-iconicity duality. Chalcogenide perovskites clearly belong to the ionic side of the spectrum, in stark contrast to conventional four-fold coordinated semiconductors. As a crude quantification of the effective ionicity/ covalency, one may calculate the Pauling electronegativity ${\left(\chi\right.}$ ) differ­ ence per anion [34]. For $\mathrm{BaZrS}{3}$ , $\begin{array}{r}{\Delta\chi=\frac{\left[3\chi(S)-\chi(\mathrm{Ba})-\chi(\mathrm{Zr})\right]}{3}=1.84\mathrm{eV}.}\end{array}$ . Using the same methodology, $\Delta\chi$ is found to be $1.43,\mathrm{eV}$ for AlN and 0.37 eV for GaAs. To put it into perspective, $\mathrm{BaZrS}{3}$ has a moderate band gap of $1.8;\mathrm{eV}$ , close to that of GaAs; but its effective ionicity is even higher than that of wide gap AlN with a band gap of $6,\mathrm{eV}$ . Possessing a moderate band gap with an exceptionally large ionicity puts $\mathrm{BaZrS}_{3}$ and other chalcogenide perovskites in a unique category of defect tolerant semi­ conductors that may offer new opportunities for electronic and opto­ electronic applications.

The high processing temperature of $\mathrm{BaZrS}{3}$ thin films in earlier studies was mainly attributed to the chemical reaction pathways with high activation barriers. As both group IIA and group IVB elements in perovskites have low electronegativity, their oxides are very stable and difficult to be converted into chalcogenides. In the approach of cosputtering of $\mathrm{BaS/Zr}$ [29], the high processing temperature above $900\ ^{\circ}\mathrm{C}$ may also be related to the difficulty of the chemical reaction between BaS and $Z\mathbf{r}$ , as $Z\mathbf{r}$ is a refractory metal with a high melting point of $1855,^{\circ}\mathrm{C}$ . The residual oxides during the handling may further exacerbate the formation of $\mathrm{BaZrS}{3}$ using sulfur. Therefore, to reduce the processing temperature, a change of the reaction pathway is necessary. PLD has been used to deposit a wide range of compounds including oxides, nitrides and carbides [35–37]. PLD is also known to be able to keep the stoichiometry of the target material under optimal conditions [38]. Postulating that by PLD deposition of a stoichiometric $\mathrm{BaZrS}{3}$ compound target under optimal conditions, $\mathrm{BaZrS}{3}$ perovskite thin films can be obtained under milder conditions.

We found that the stoichiometry of the Ba:Zr cations of the asdeposited films is dependent on laser fluence of the PLD system. This effect can be understood as primarily due to the large difference in the cohesive energy of Ba vs. Zr (1.9 eV/atom vs. 6.25 eV/atom). Therefore, low laser fluence is expected to result in Ba-rich films, while very high laser fluence may result in Ba deficiency due to reverse sputtering. Under a narrow window of fluence of close to $5~\mathrm{J/cm^{2}}$ , a 1:1 ratio of Ba: Zr was obtained, as shown in Fig S2(a), However, it was found that the as-deposited films were amorphous, even with an in-situ substrate temperature as high as $700\ {}^{\circ}\mathrm{C}$ (Fig S3). It was further determined that these films are highly sulfur deficient, with a composition of $\mathrm{{BaZrS}_{2.4},}$ as shown by the EDX spectroscopy results (Fig S2(b)), which perhaps contributes to their poor crystallinity.

To obtain stoichiometric, single phase $\mathrm{BaZrS}{3}$ in the orthorhombic perovskite structure, thermal annealing of the as-deposited amorphous $\mathrm{BaZrS{2.4}}$ films in $\mathsf{C}\mathsf{S}{2}$ at various temperatures ranging from $500^{\circ}\mathrm{C}$ to $900^{\circ}\mathrm{C}$ were performed in a two-zone tube furnace. All of the $\mathrm{BaZrS}{3}$ films after annealing appeared dark red in color (see optical images in Fig. 1(b)), except the one annealed at $500\ ^{\circ}\mathrm{C}{:}$ , indicating strong visible light absorption. The $500\ ^{\circ}\mathrm{C}$ sample does show lighter color, suggesting incomplete crystallization. The sulfur concentration as measured by EDX is annealing temperature dependent, as seen in Fig S4. While asdeposited $\mathtt{B a Z r S}{2.4}$ films are sulfur deficient, the composition of the films progressively approaches perfect stoichiometry of $\mathrm{Ba}{:}\mathrm{Zr}{:}\mathrm{S}{=1{:}1{:}3}$ with decreasing annealing temperature. To investigate the annealing temperature dependent structural evolution of the $\mathrm{BaZrS}{3}$ films, XRD θ-2θ scans were performed for films annealed at $900^{\circ}\mathrm{C}$ , $850^{\circ}\mathrm{C}$ , $650\ ^{\circ}\mathrm{C}{\mathrm{:}}$ , $550^{\circ}\mathrm{C}$ , and $500^{\circ}\mathrm{C}$ , respectively, as shown in Fig. 1(a). The peaks labeled with Miller indices indicate standard peak positions of $\mathrm{BaZrS}_{3}$ with a Pnma structure, retrieved from the PDF card JCPDS 00-015-0327. It is found that all discernible peaks for the samples annealed from $500\ ^{\circ}\mathrm{C}$ to $900\ {}^{\circ}\mathrm{C}$ match well with the reference. No additional peaks were observed, suggesting that the obtained films were of a single phase. Some diffraction peaks with low intensities were not clearly observed from films annealed at $550^{\circ}\mathrm{C}$ and $500^{\circ}\mathrm{C}$ , as a result of low crystallinity and small grain sizes at such low temperatures. As can be seen from

Fig. 1(b), the full width at half maximum (FWHM) of the (121) diffraction peak decreases from $0.62^{\circ}$ to $0.45^{\circ}$ with increasing annealing temperatures from $550^{\circ}\mathrm{C}$ to $900^{\circ}\mathrm{C}$ . The crystallite sizes as calculated from Scherrer’s formula using the FWHM of the (121) diffraction peak decreases monotonically with decreasing annealing temperature. The lateral grain sizes, as imaged by AFM and SEM (shown in Fig. 2(a)–(h) and Fig S5) also decreases monotonically from ${\sim}110;\mathrm{nm}$ at $900^{\circ}\mathrm{C}$ to $35;\mathrm{nm}$ at $550^{\circ}\mathrm{C}$ . It is interesting to note that the lateral grain sizes are several times larger than the vertical grain sizes. Understanding the detailed crystallization mechanisms in the future will help to obtain films with enhanced grain size and crystallinity.

The surface roughness of a semiconductor thin films plays an important role in determining the interface quality, carrier transport in junctions, and thus device performance [39,40]. It is expected that a lower annealing temperature with smaller crystallite size should also possess lower surface roughness. The AFM top-view images in Fig. 2(a)– (d) and 3D topography images displayed in Fig. 2(e)–(h) show clearly that the crystallite size and surface roughness of the $\mathrm{BaZrS}{3}$ thin films decreases consistently with decreasing annealing temperature from $900^{\circ}\mathrm{C}$ to $550^{\circ}\mathrm{C}$ , respectively. The height profile measured along a scan length of $5,\upmu\mathrm{m}$ for different annealing temperatures is shown in Fig. 2(i). The root mean square (RMS) of the surface roughness $\mathbf{(R{q})}$ decreases by 20 folds from $10.5;\mathrm{nm}$ at annealing temperature of $900~^{\circ}\mathrm{C}$ to $0.59;\mathrm{nm}$ at annealing temperature of $650\ ^{\circ}\mathrm{C}_{\mathrm{:}}$ , as seen in Fig. 2(j). This trend can be explained by the consequence of a decrease in grain size with decreasing annealing temperature.

The room temperature Raman spectra of $\mathrm{BaZrS}{3}$ films annealed at $900^{\circ}\mathrm{C}$ , $850^{\circ}\mathrm{C}$ , $650^{\circ}\mathrm{C}$ , and $550^{\circ}\mathrm{C}$ are shown in Fig. 3(a). Several peaks can be identified in the range of $50;\mathrm{cm}^{-1}$ to $500;\mathrm{cm}^{-1}$ and six of them are assigned to be $B{1g}^{1},\ A_{g}^{4},\ B_{2g}^{6},\ B_{1g}^{4},\ B_{1g}^{5}$ and $B_{3g}^{5}$ vibrational modes calculated theoretically for $\mathrm{BaZrS}{3}$ with the Pnma structure [27]. The Raman spectra measured at low temperature (100 K) are provided in Supporting Information Fig. S6. In addition, we measured the Raman spectra at higher wavenumbers between $1000\mathrm{cm}^{-1}$ and $2000\mathrm{cm.}^{1}$ for all the films to investigate the possibility of carbon contamination resulting from $C S{2}$ decomposition. As can be seen from Fig. 3(b), for the film annealed at $900\ ^{\circ}{\mathrm{C}},$ , two peaks at 1330 and $1620~\mathrm{cm}^{-1}$ were observed, which can be attributed to the D and G bands of carbon ma­ terials such as graphite and nano-diamonds. While these modes are barely observable for the sample annealed at $850\ ^{\circ}\mathrm{C}{\mathrm{:}}$ , they are absent in the films annealed at and below $650;^{\circ}\mathrm{C}$ . This suggests that carbon contamination has been eliminated by low temperature processing. The presence of carbon also explains the color difference of $\mathtt{B a Z r S}{3}$ processed at different temperatures: while the films processed at $1050\ ^{\circ}0$ show black color (not shown), the films processed at $900\ {}^{\circ}\mathrm{C}$ and below are dark red.

Fig. 1. (a) XRD patterns of $\mathrm{\bfBaZrS_{3}}$ films annealed at different temperatures. (b) The FWHMs of the (121) peak as a function of annealing temperature. The inset shows the optical images of the $\mathrm{BaZrS_{3}}$ films annealed at the respective temperatures.

Fig. 2. AFM top-view and 3D topographic images of $\mathrm{BaZrS_{3}}$ thin films annealed at (a) and (e) $900;^{\circ}\mathrm{C},$ (b) and (f) $850;^{\circ}\mathrm{C}{:}$ , (c) and (g) $650;^{\circ}\mathrm{C}$ , and (d) and (h) $550~^{\circ}\mathrm{C}$ . (i) The height profile measured along a scan length of $5,\upmu\mathrm{m}$ . (j) The root mean square (RMS) surface roughness $\mathbf{(R{q})}$ as a function of annealing temperature.

$\mathrm{BaZrS}{3}$ thin films annealed at $650\ ^{\circ}\mathrm{C}$ and $900\ ^{\circ}\mathrm{C}$ were selected for UV–Vis absorption measurements to examine their optical properties. The thicknesses of these films are about $280;\mathrm{nm}{\mathrm{,}}$ , as measured by AFM (Fig. 4(a)). The absorption spectra are shown in Fig. 4(b). It can be seen that their absorption coefficient $\upalpha$ above the band gap are $\sim5\times10^{4}$ and $\sim6\times10^{4},\mathrm{cm}^{-1}$ , respectively. These values are somewhat smaller than that of $\mathrm{BaZrS}{3}$ films processed at higher temperatures [28]. A linear fitting near the band edge of the Tauc plot results in a band gap value of 1.98 and $1.86;\mathrm{eV}$ , respectively, for the films annealed at $650\ ^{\circ}\mathrm{C}$ and $900\ {}^{\circ}\mathrm{C}{:}$ , respectively, as shown in Fig. 4(c). These values are slightly higher than the band gap in the films processed at $1050;^{\circ}\mathrm{C}$ . Several factors could contribute to these differences: 1. The presence of a minute amount of amorphous oxides/oxysulfides, possibly introduced in the target processing, could not be completely eliminated in low tempera­ ture processing. The wide gap oxides/oxysulfides may lead to an over­ estimate of the band gap, as well as lower the absorption coefficient. Careful examination of XRD patterns of samples processed at different temperatures (Fig. 1(a)) reveals no oxide peaks or shift of the peaks compared to the $\mathrm{BaZrS}{3}$ standard, within the instrument limit. This rules out the presence of measurable amount of crystalline oxides or oxy­ sulfides; 2. The presence of carbon contamination in previous high temperature processed samples may lead to overestimate of the ab­ sorption coefficient and underestimate of band gap, as carbon absorbs in a broad range of wavelengths; 3. We observed a small but noticeable shift of the Raman peaks to higher wavenumbers at higher processing temperatures (Fig. S6). While presence of oxysulfides could lead to shift of Raman modes to higher frequencies, this is inconsistent with the fact that this shift is for samples processed at higher temperatures and therefore conversion to the $\mathrm{BaZrS}{3}$ phase is more complete. We suggest that a slight difference in structural distortion (octahedron rotation) at different processing temperatures is responsible for the observed Raman shift, which could contribute to subtle differences in electronic struc­ tures and band gaps [41].

In summary, lower temperature growth of $\mathrm{BaZrS}{3}$ results in stoi­ chiometric films with minimized concentration of sulfur vacancies, eliminated carbon contamination, smaller crystallite sizes and lower surface roughness. These are expected to strongly affect the carrier transport characteristics and device performance. To investigate the influence of film microstructure on carrier transport, we fabricated and measured photodetector and FET devices using the $\mathrm{BaZrS}{3}$ films annealed at different temperatures as active semiconductors.

The I-V curves of the photodetector were measured in the dark and under illumination, for films annealed at temperatures of $850^{\circ}\mathrm{C}$ $650^{\circ}\mathrm{C}$ , and $550^{\circ}\mathrm{C}$ , respectively, as shown in Fig. 5(a)–(c). An optical image of the device is shown in the inset of Fig. 5(a). At $850^{\circ}\mathrm{C}$ , the I-V curves show little difference under dark and illumination. This suggests that carriers available for the transport are essentially intrinsic to the material, and the contribution from the photo-generated carriers to the transport is negligible. The high dark current can be attributed to the high concentration of sulfur vacancies, as each sulfur vacancy can donate two electrons. Using the sulfur vacancy concentration measured by EDX (Fig S4), we estimate the carrier density to be $1.7\times10^{21},\mathrm{cm}^{-3}$ For the film annealed at $650^{\circ}\mathrm{C}$ , the measured photocurrent is 80 times greater than the dark current, as can be seen in Fig. 5(b). The film annealed at $550^{\circ}\mathrm{C}$ shows the highest photocurrent of $17,\tt n A$ at $_{10\mathrm{V}}$ However, the on/off ratio is only 20 times due to the relatively higher dark current. The relatively higher dark current could be caused by the lower crystallinity, where a number of defects may be present and contribute to excess carriers, as shown by our unpublished theoretical study [42]. The photo responsivity is calculated to be $0.08;\mathrm{mA/W}$ at a bias voltage of $5,\mathrm{V}$ for the films annealed at $550^{\circ}\mathrm{C}$ . The moderate photo response is primarily due to the relatively small grain sizes, which leads to significant carrier scattering and limits the carrier mobility. In Fig. 5 (d), we plot the photocurrent as a function of time for the films annealed at $650^{\circ}\mathrm{C}$ under periodically switched light source at a fixed bias voltage of $8,\mathrm{V}$ . The device shows fast response and low rise/fall time. Due to instrument limit, we estimate the upper limit of the response time to be $0.3\mathrm{ms}$ . However, the actual response time could be much smaller. The I-V curves measured at different excitation wavelengths for films annealed at $650^{\circ}\mathrm{C}$ are plotted in Fig. 5(e), and the wavelength depen­ dent photocurrent at bias voltages of 3, 5 and $8,\mathrm{V}$ are shown in Fig. 5(f).

Fig. 3. (a) The room temperature Raman spectra of $\mathrm{BaZrS}_{3}$ films annealed at different temperatures. (b) Raman spectra at higher wavenumbers showing presence/ absence of carbon contamination.

Fig. 4. (a) The AFM height profiles, (b) The absorption coefficient, and (c) The Tauc plot of the films annealed at $650;^{\circ}\mathrm{C}$ and $900~^{\circ}\mathrm{C}$ , respectively.

The device shows the highest photo response at the excitation wave­ length around $500;\mathrm{nm}$ . The cutoff wavelength is found to be about $670;\mathrm{nm}$ . This confirms the band gap of the $\mathrm{BaZrS}_{3}$ film annealed at $650~^{\circ}\mathrm{C}$ to be about $1.85;\mathrm{eV}$ . This suggests that the UV–vis absorption measurements slightly overestimate the band gap, perhaps due to the small amount of oxides.

FET devices were fabricated using the film annealed at $650^{\circ}\mathrm{C}$ to further study the carrier transport, using an IL side gate, as shown by the optical image in the inset of Fig. 6(a). The drain-source current $\mathrm{I_{DS}}$ as a function of bias voltage $\mathrm{{V_{DS}}}$ , measured at positive and negative gate voltages $\mathrm{V}_{\mathrm{G}}$ are shown in Fig. 6(a) and (b), respectively. The total spe­ cific capacitance of the $\mathrm{IL}$ is determined by the capacitance of the electrical double layer (EDL), and can reach as high as ${\sim}20\upmu\mathrm{F/cm^{2}}$ [43], which makes it possible to generate high electric field of $\sim$ $10\ \mathrm{MV/cm}$ at low gate voltages [44], comparing to commonly used metal oxide dielectric such as $\mathrm{{SiO}{2}.}$ , $\mathrm{{Al}}{2}0_{3}$ or $\operatorname{HfO}{2}$ . At a $\mathrm{V{DS}}$ of $1;\mathrm{V}$ , IDS is measured to be $3\times10^{-11}$ A without gating, which suggests very small intrinsic carrier density. The intrinsic carrier type is found to be n-type, which is consistent with our previous work [28]. The intrinsic carrier concentration is found to be $4.4\times10^{10}~\mathrm{{cm}}^{-3}$ . These results further confirm that carrier concentration is suppressed by reducing the density of sulfur vacancies in $\mathrm{BaZrS}{3}$ films processed at low temperatures. With increasing $\mathrm{V{G}}$ to ${1\textrm{V}}$ , IDS reaches to $2\times10^{-7},\mathrm{A},$ resulting in an ON/OFF ratio of $6.5\times10^{3}$ . Similar but smaller $\mathrm{I{DS}}$ can be observed for negative $\mathrm{V_{G}}$ , as shown in Fig. 6(b), suggesting ambipolar behavior in our FETs. We further estimate the electron and hole mobilities to be $16.8;\mathrm{cm}^{2}/\mathrm{Vs}$ and $2.6;\mathrm{cm}^{2}/\mathrm{Vs}$ , respectively, from the saturated $\mathrm{I_{DS}}$ vs. $\mathrm{V}{\mathrm{G}}$ data. The capacitance of the IL layer used in the calculation was obtained from the C-V measurements. The carrier mobility is primarily limited by the small grain sizes which contribute to substantial carrier scattering. Future work should focus on improving the grain size of $\mathrm{BaZrS}{3}$ films processed at low temperatures to improve the performance of electronic and op­ toelectronic devices.

4. Conclusion

In conclusion, we have obtained polycrystalline perovskite $\mathrm{BaZrS}{3}$ thin films, by low temperature crystallization of PLD deposited amor­ phous $\mathtt{B a Z r S{x}}$ films in $\mathrm{C}S_{2}$ . The films with good crystallinity and a sur­ face roughness smaller than $0.6;\mathrm{nm}$ can be obtained at processing temperatures as low as $550~^{\circ}\mathrm{C}$ . The low processing temperatures result in the elimination of both excessive carriers due to vacancy formation and carbon contamination. Prototype photodetector devices fabricated with such films show good visible light sensitivity, and their wavelength dependent photo response confirms a band gap of ${\sim}1.85,{\mathrm{{eV}}}$ . The FET indicates ambipolarity of the charge carriers, with electron and hole mobilities of $16.8;\mathrm{cm}^{2}/\mathrm{Vs}$ and $2.6;\mathrm{cm}^{2}/\mathrm{Vs}$ , respectively. Our findings will expedite the progress towards applications of $\mathrm{BaZrS}_{3}$ and other chalcogenide perovskites for electronics and optoelectronics.

CRediT authorship contribution statement

Zhonghai Yu: Investigation, Validation, Methodology, Formal analysis, Writing - review & editing, Visualization. Xiucheng Wei:

Fig. 5. The I-V curves in dark and under illumination of the photodetectors fabricated from the films annealed at (a) $850^{\circ}\mathrm{C}$ , (b) $650;^{\circ}\mathrm{C}$ , and (c) $550^{\circ}\mathrm{C}$ . The inset in (a) is an optical image of the device. (d) Current vs time for the device using the film annealed at $650;^{\circ}\mathrm{C}$ , with a periodically switched light source $\left(532,\mathrm{nm}\right)$ at the fixed bias of $8,\mathrm{V}$ (e) The I-V curves measured at different excitation wavelengths for the device in (d). (f) The wavelength dependent photocurrent at bias voltages of 3, 5 and $8,\mathrm{V}$ extracted from (e).

Fig. 6. The $\mathrm{I_{DS}}$ vs $\mathrm{V_{DS}}$ of the FET fabricated using the film annealed at $650;^{\circ}\mathrm{C}$ under (a) positive and (b) negative gate voltages $\mathrm{V_{G}}$ . The inset in (a) is an optical image of the FET with IL side-gate.

Investigation, Methodology, Validation, Formal analysis, Writing - original draft, Writing - review & editing, Visualization. Yixiong Zheng: Investigation, Validation, Formal analysis. Haolei Hui: Investigation. Mengying Bian: Investigation. Samyak Dhole: Investigation. JungHun Seo: Resources, Writing - review & editing, Supervision. Yi-Yang Sun: Methodology. Quanxi Jia: Methodology, Writing - review & editing, Funding acquisition. Shengbai Zhang: Methodology. Sen Yang: Funding acquisition, Supervision. Hao Zeng: Conceptualization, Validation, Methodology, Formal analysis, Writing - original draft, Writing - review & editing, Supervision, Funding acquisition.

Declaration of Competing Interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

Acknowledgements

Work supported by United States National Science Foundation (US NSF) CBET-1510121, CBET-1510948, MRI- 1229208, and Department of Energy (DOE) DE-EE0007364. Z. Yu and S. Yang acknowledge Na­ tional Natural Science Foundation of China (NSFC) (Grant No. 91963111) and Key Scientific and Technological Innovation Team of Shaanxi Province (2020TD-001). Y. -Y.S. acknowledges support by NSFC under Grant No. 11774365.

Author contributions

Z.Y. synthesized the PLD target and thin films, and performed structural and optical characterizations. X.W. fabricated the photode­ tector and FET devices. X.W. and Y.Z. performed device characteriza­ tions and data analysis. H. Z. conceived the idea. S.Y., J.S. and H.Z.

supervised the project. X.W., Z.Y. and H.Z. wrote the manuscript. All authors discussed the results and contributed to the final manuscript.

Appendix A. Supporting information

Supplementary data associated with this article can be found in the online version at doi:10.1016/j.nanoen.2021.105959.

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Zhonghai Yu received his bachelor of science degree from Taiyuan University of Technology in 2014. He is currently a doctoral student working on his Ph.D., focusing on Chalco­ genide Perovskites solar cell at Xi’an Jiaotong University since 2015. He will be a visiting Ph.D. student at the University at Buffalo.

Xiucheng Wei received his B.S. degree in physics from Xi’an Jiaotong University, China. He is currently a graduate student in the department of physics at University at Buffalo. His research interests focus mainly on novel photovoltaic semi­ conducting materials and devices.

Dr. Jung-Hun Seo received his B.S. degree in electronics and electrical engineering from Korea University, Seoul, Republic of Korea, in 2006. He received his M.S. and Ph.D. degrees in Electrical and Computer Engineering from University of Wisconsin-Madison in 2011 and 2014, respectively and worked as an assistant scientist at the University of Wisconsin-Madison from 2014 to 2016. Since 2016, he has been an assistant pro­ fessor at the Department of Materials Design and Innovation, University at Buffalo. His research interests mainly focus on the synthesis and heterogenous integration of low dimensional wide bandgap semiconductors and their high frequency and high power devices and flexible device applications.

Yixiong Zheng is currently pursuing his Ph.D. degree in Prof. Jung-Hun Seo’s group, at the Department of Materials Design and Innovation, University at Buffalo. He received his B.S. degree in optoelectronics engineering from University of Electronic Science and Technology of China in 2014. His research interest is the design and fabrication of wide band gap semiconductor devices.

Yi-Yang Sun received his Ph.D. in physics from National University of Singapore (NUS) in 2004. Since then, he has worked as a postdoc at NUS, National Renewable Energy Laboratory and Rensselaer Polytechnic Institute (RPI), USA. In 2010, he was appointed Research Assistant Professor and later Research Scientist at RPI. In 2017, he assumed a Professor position at Shanghai Institute of Ceramics, Chinese Academy of Sciences. He has been working on first-principles study of energy-related materials.

Haolei Hui received his bachelor’s degree in Xi’an Jiaotong University in 2018. He is currently a Ph.D. student in depart­ ment of physics at University at Buffalo. He is focusing on novel chalcogenide perovskite as photovoltaic materials.

Quanxi Jia is an Empire Innovation Professor and National Grid Professor of Materials Research at the University at Buf­ falo (UB). Prior to joining UB in 2016, he had worked at Los Alamos National Laboratory for 24 years, with the last two years serving as the co-Director and then Director of the Center for Integrated Nanotechnologies, a US Department of Energy Nanoscale Science Research Center operated jointly by Los Alamos and Sandia National Laboratories. His research focuses on nanostructured and multifunctional materials, with a particular effort on the synthesis and study of processingstructure–property relationships of epitaxial films for energy and electronic applications.

Mengying Bian received her Ph.D. in Physics from the Peking University in 2018. Currently, she is working as a postdoctoral researcher in the group of Professor Yanglong Hou at Peking University and a visiting scholar in the group of Professor Hao Zeng at University at Buffalo, where her research is focused on 2D magnetic materials and devices.

Shengbai Zhang received his Ph.D. in physics from University of California at Berkeley in 1989. He moved to Xerox PARC as a postdoc, before joining the National Renewable Energy Labo­ ratory in 1991. In 2008, he became the Senior Kodosky Constellation Chair and Professor in Physics at Rensselaer Polytechnic Institute. His expertise is first-principles theory, modeling, and calculation. Recent work involves chalcogenide perovskites for photovoltaic, ultrafast phase change memory materials, topological carbon networks, non-equilibrium growth, excited state dynamics, and unconventional twodimensional materials and excitonic insulators. He is a Fellow of the American Physical Society since 2001.

Samyak Dhole received his M.Sc. in Chemistry from BITS Pilani, India in 2018. He is currently a graduate student in the Department of Materials Design and Innovation in University at Buffalo. His research interests lie primarily in thin films, synthesis and characterization of perovskite-based materials, resistive switching and polymer assisted deposition.

Sen Yang received his Ph.D. in materials physics from the Xi’an Jiaotong University (XJTU), China in 2005. He joined the National Institute for Materials Science, Japan in 2005 as a JSPS (Japan Society for the Promotion of Science) post-doctor. In the year of 2010, he came back to XJTU and was promoted to full professor in 2013. His research interests are in magnetism and magnetic materials, smart materials, phase transition and so on.

Hao Zeng received his Ph.D. in physics from the University of Nebraska-Lincoln in 2001. He joined the Department of Physics at the University at Buffalo in 2004 after a 3-year postdoc at IBM T.J. Watson Research Center, and was promoted to full professor in 2014. His research interests are in nanoscale magnetism and magnetic materials, spintronics, 2D materials and unconventional semiconductor materials and devices.