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Thin-Film Deposition of Surface Passivated Black Phosphorus Nezhueyotl Izquierdo,† Jason C. Myers,‡ Nicholas C. A. Seaton,‡ Sushil K. Pandey,§ and Stephen A. Campbell*,§
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†
Department of Chemical Engineering and Materials Science, University of Minnesota, Minneapolis, Minnesota 55455, United States ‡ Characterization Facility, University of Minnesota, Minneapolis, Minnesota 55455, United States § Department of Electrical and Computer Engineering, University of Minnesota, Minneapolis, Minnesota 55455, United States S Supporting Information *
ABSTRACT: A single-step, direct silicon-substrate growth of black phosphorus (BP) crystals is achieved in a self-contained short-way transport technique under low-pressure conditions (102. KEYWORDS: two-dimensional, black phosphorus, nucleation, growth, thin film
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lack of a band gap prevents most applications in traditional semiconductor devices, particularly for FET switching that requires a high on/off ratio and low off-state current.6,7 This graphene device application limitation increased attention on other suitable 2D semiconductor material options, such as the transition-metal dichalcogenides (TMD)8,9 which have a semiconducting bandgap, typically in the range of 0.8−2.0 eV.8 This is ideal for making FETs and optoelectronics. There are hundreds of known TMD materials and likely many more that remain undiscovered. However, the carrier mobility of these materials is generally in the 10−200 cm2/v·s range, well below that of graphene.10 BP contains a direct band gap similar to those of the TMDs, ranging from 0.3 in the bulk to about 1.9 eV for monolayer BP which is commonly known as phosphorene. As with the TMDs, this allows for the electrical properties of the material
ver the past decade, the two-dimensional (2D) material class has been mined for materials with unique and phenomenal properties as the foundation for a potential post-silicon era.1 The miniaturization of semiconductor device sizes in an effort to increase chip density requires reduction in all three dimensions, which has led research into single- or few-layer-thick materials.2 However, decreasing the thickness of 3D semiconductors reduces the mobility by a factor of t6 (t = thickness), due to charge carriers scattering caused by dangling bonds.1 Two-dimensional materials have an extremely high surface area to volume ratio and no dangling bonds in the vertical direction.3 Instead, layers are bound by van der Waals forces. As a result, 2D materials do not face the issue with scaling the thickness, even at the atomic scale. Other advantages of 2D material include limiting the cost of the material required, the potential for clean heterostructure interfaces without an ultrahigh-vacuum growth process, and a layer-dependent bandgap.4,5 Graphene was initially at the forefront of 2D research, considering its record charge carrier mobility of 106 cm2/v·s. However, graphene is a semimetal. Its © 2019 American Chemical Society
Received: March 27, 2019 Accepted: May 30, 2019 Published: May 30, 2019 7091
DOI: 10.1021/acsnano.9b02385 ACS Nano 2019, 13, 7091−7099
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Cite This: ACS Nano 2019, 13, 7091−7099
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small needles (Supplementary Figure 1). Efforts to characterize the needles were unsuccessful due to laser and electron beam ablation. Some BP did form, but at the cold end of the ampule, not on the substrate. This is unsurprising since the thermal gradient is designed to separate phases and localize crystal growth to a specific site in order to grow large crystals. To deposit a uniform film over a large area, one should keep the ampule at a uniform temperature. BP microribbons have previously been synthesized by Zhao et al. by removing the thermal gradient, however, the starting material weight of tin, iodine, and phosphorus was not published.31 BP growth parameters were investigated by keeping the phosphorus mass (500 mg) constant and altering the phosphorus to Sn and SnI4 ratio, as depicted in Figure 1.
to be tailored to specific devices, such as optoelectronics and transistors.11,12 Historically, the synthesis of bulk black phosphorus (b-BP) has required pressure in excess of 1 GPa. The field has received considerable scientific attention since Nilges et al. introduced a low-pressure, short-way transport method for the synthesis of b-BP using red phosphorus, tin, and iodine in the form of SnI4 (RP-Sn-I).13,14 The process produces crystals that are multiple millimeters long inside a fused silica ampule. We and others have replicated the process with excellent results. A major problem for phosphorene development revolves around the inability of either liquid or mechanical exfoliation methods to create flakes that can cover large areas with welldefined orientation and thickness. Obtaining monolayer flakes (phosphorene) is particularly difficult. As a result, device fabrication requires direct writing techniques, has reproducibility issues, and is limited to small discrete devices.15−17 Ultimately, a synthesis method to grow large-area phosphorene films under easily obtained conditions is required in order for BP commercialization. A reasonable first step toward this goal would be the direct growth of BP thin films onto substrates. This has been studied by two groups who both produced phosphorene as well as b-BP, simultaneously. Li et al. coated red phosphorus on a sapphire substrate, transferred hexagonal boron nitride to cover the red phosphorus islands, and used a heated piston cylinder to apply 700 °C and 1.5 GPa, forming a BP film.18 Smith et al. coated a silicon substrate with red phosphorus at 600 °C, followed by a pressure vessel reactor treatment with a pressure reaching 6.895 MPa and a temperature of 950 °C.19 Thus, both of these BP growth methods require a very high-pressure environment and/or very high temperatures. High-temperature and pressure growth conditions limit accessible growth substrates and suitable complementary heterostructure device material. Phosphorene device research and development would benefit from a growth method that occurs under milder growth conditions. A reasonable route to achieving direct growth of BP films is to start from the Nilges b-BP method. Their RP-Sn-I system has the most literature available as well as the highest red phosphorus to BP conversion rate (95%).20 Unfortunately, the growth mechanism of BP using the RP-Sn-I system has not yet been identified, making process adaptations difficult. It was recently proposed that violet (monoclinic or Hittorf’s) phosphorus (VP) is an intermediate phase in the transformation of red phosphorus to BP using this growth system.21 VP, a wide bandgap and high-mobility material, has a 2D layered crystal structure with van der Waals interlayer bonds, but a covalent phosphorus polymeric unit cell.22−25 A second important limitation of BP is its susceptibility for degradation due to exposure to H2O vapor, O2, and UV blue light.26,27 Phosphorus oxide forms on the surface of exposed layers and dissolves in H2O from the air.28,29 Thus, current device research requires making devices as rapidly as possible once the flake has been exfoliated from the b-BP crystal. It is customary to coat a thin passivation layer, such as Al2O3 or Si3N4, when fabricating devices to protect the film from environmental degradation.28,30 If a thin-film BP is to be useful, it must also avoid this degradation effect since the reason for thin-film deposition is to avoid exfoliation.
Figure 1. Successful growth of BP on a silicon wafer is designated by the black circle. The optimal growth condition is represented by the black star. Red data points indicate samples that did not produce any BP on the silicon wafer. The purple diamond is the standard ratios used for bulk-BP growth. SnI4/P and Sn/P weighted ratios are calculated based on 500 mg of phosphorus starting material.
Each of the black (circle) and violet (circle) data points in Figure 1 represents a successful reaction parameter for synthesizing BP and VP, respectively, directly on a silicon substrate. The red (circle) data points indicate samples where no BP was detected on the surface of the wafer. The purple diamond (Figure 1) indicates the standard Nilges b-BP process mass ratio of 25:2:1 of phosphorus, Sn, and SnI4 but without a thermal gradient.13 Under this condition, black amorphous phosphorus nodules were primarily formed when the thermal gradient was removed in the presence of a silicon wafer. Increasing the amount of iodine in the form of SnI4 in the reactor quickly brought the reaction into a regime that formed thin-film BP crystals on the surface of the wafer. The behavior persisted for a broad range of iodine concentrations. Conversely, decreasing the SnI4 concentration resulted in the formation of VP, in both 2D and bulk forms. Previous VP growth methods involved far more difficult processing including prolonged annealing periods (5−7 days), plasmaassisted synthesis, or highly reactive and pyrophoric Group 1 alkali metals, such as rubidium and cesium.22,23,32 A link between VP and BP has also been proposed. Smith et al. suggested that VP may be a transition phase for BP synthesis in a Sn:I:P growth system.19 Optical microscope imaging of VP shows the 2D nature of the material evident from the steps and light interference patterns (Figure 2a). Raman confocal mapping of VP shows
RESULTS Attempts to deposit BP on Si using the standard Nilges recipe resulted in a large amorphous phosphorus mass and some 7092
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Figure 2. (a) Optical microscope image of VP. (b) Raman map of VP. (c) Optical microscope images of BP. (d) Raman map of BP.
that the intensity of the phonon mode at 356 cm−1 is independent of position (Figure 2b). Thus, the VP is crystalline along the entire surface. Raman spectroscopy of VP shows characteristic intratube modes above 340 cm−1, bond angle distortions, rotational vibration, and bending modes below 310 cm−1(Supplementary Figure 2).33 The peak observed at 520 cm−1 is attributed to the substrate silicon. The range of BP thickness calculated from 12 flakes is 30−450 nm. The optical microscope image of the hero BP microbelt (Figure 2c) shows a uniform crystal surface with dimensions of 11 × 81.6 μm. Confocal Raman mapping of BP B2g phonon shows a uniform and highly crystalline material (Figure 2d). Ag1, B2g, and Ag2 characteristic BP phonon modes are found at 362.2, 438.5, and 466.1 cm−1, respectively (Supplementary Figure 2a). The BP flake is 117 nm thick according to an atomic force microscopy (AFM) measurement (Supplementary Figure 4). Using Auger electron spectroscopy (AES), we were able to find a layer of Sn on the surface of the flake. The AESmeasured Sn atomic percentage, AES depth penetration (5−10 nm), and Sn atomic radius (217 ppm) were used to estimate the Sn film thickness as 2.0 ± 0.5 nm. This Sn surface layer functions as an in situ passivation layer that greatly inhibits degradation to exposure to ambient conditions. We found that b-BP grown by the Nilges method also has a thin film of Sn on the surface. However, the commonly used mechanical exfoliation process exposes unprotected BP surfaces which rapidly degrade under ambient conditions.26,27,34 Thus, passivation occurs in situ under growth conditions without the presence of known oxidizing agents, such as O2 and H2O, and can later be removed during device fabrication. Previous published methods to synthesize BP directly onto substrates have not shown increased stability or resistance to degradation under ambient conditions. An AES surface scan indicates the presence of phosphorus, oxygen, carbon, iodine, and tin on b-
BP that we grew using the Nilges method. After the first sputter step (0.5 min) however, the only element that remains evident (greater than about 1%) is phosphorus. The AES mapping of b-BP prior to the AES sputtering indicates a thin, uniform coverage of oxygen, phosphorus, and tin. AES analysis does not provide chemical bonding information. Therefore, X-ray photoelectron spectroscopy (XPS) was performed on the surface of b-BP. XPS analysis of the b-BP surface provides a close model for our single-step, direct substrate growth of BP flakes and surface chemistry. XPS of a new BP sample is shown in Supplementary Figure 6. The primary peak at 130 eV and the shoulder peak at 131 eV are characteristic of P3/2 and P1/2, respectively, indicating that phosphorus in the as-grown sample is primarily in the elemental form.35 A weak phosphorus oxide peak present at ∼135 eV in the as-grown sample is believed to be due to ambient exposure and resultant oxidation during sample transport and preparation (approximately 1 h).35 The phosphorus oxide is consistent with an oxidation state of P5+.36 A significant increase in the phosphorus oxide (P5+) peak is seen in the oxidized sample (Supplementary Figure 6b) after 48 h of room-temperature exposure to ambient conditions, as expected. The XPS measurement for oxygen reveals adsorbed oxygen at 532 eV for the as-grown sample, followed by a shift in the oxidized sample to 533 eV as a result of phosphorus oxidation.37 Our XPS analysis is consistent with the reported phosphorus oxidation compound phosphorus pentoxide, P2O5, as the BP oxidation compound that arises from prolonged exposure to ambient conditions.35,38 Oxidation of Sn is not observed in XPS data. Furthermore, the XPS measurement for Sn 3d5/2 (488 eV) and 3d3/2 (497 eV) is consistent with elemental Sn.39 Carbon (C1) and tin (Sn 5d) peaks are unaffected after exposure to ambient conditions. BP flakes that are covered by a uniform elemental Sn thin film are highly resistant to degradation under ambient 7093
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Figure 3. EBSD of a single crystal BP flake. (a) Scanning electron microscopy with 2.00 kV accelerating voltage. The yellow dashed overlay is the portion of the single crystal used for CS-TEM measurements. (b) EBSD orientation map with inverse pole figure (color legend overlay) showing the BP crystal from (a) has a unidirectional orientation.
Figure 4. (a) Cross-sectional TEM HAADF image of BP along the armchair crystallographic axis. (b) Cross-sectional TEM image of SiO2, Sn, BP, and Sn stack including a SnIx crystal inclusion.
these substrates was demonstrated to achieve BP growth using the optimal conditions for growth on silicon. A thorough optimization was not conducted to determine the best growth conditions for substrates other than bare silicon. The crystallinity of the BP grown on silicon varies from flakes with high-angle grain boundaries (>15° lattice orientation shift), low-angle grain boundaries (102 for both thin-film and bulk recipes (Figure 5c). This inversion layer mobility and on/off behavior for both carrier types are consistent with similar thickness BP devices in the literature.40 However, the I−V linear plot (inset) shows a −16 V and −12 V threshold voltage for bulk and thinfilm devices, respectively. The calculated VT shift is 4 V. This lateral shift in the curves suggests a difference in the charge in the gate oxide or a difference in the carrier concentration in the samples. The latter can be a result of a number of factors including the effects of dopants or defects in the BP related to differences in the growth process, adsorbed contaminants at the surface, and defects caused by argon plasma thinning.41
DISCUSSION While the growth mechanism of b-BP in the Nilges process is not understood, several observations have been made in the literature for bulk growth. In situ neutron diffraction measurements show the appearance of BP peaks during temperature ramp down at 500−400 C.13 The presence of iodine seems critical. For example, replacing tin iodide with tin chloride prevents the formation of b-BP.20 In contrast, b-BP is still produced when Sn is replaced with several phosphorus-soluble metals, such as In, Cd, and Pb in the presence of SnI4.20 Transport of the Sn from the hot end to the cold end of the ampule under b-BP growth conditions is likely to be one important role of the iodine.31 The equilibrium vapor pressure calculated (Supplementary Figure 7) from the Antoine equation using NIST parameters indicate that the Sn flux is extremely low compared to that of phosphorus and iodine. As a result, Sn placed at the hot end of the ampule in the presence of only phosphorus will not transport to the cold end in the 14 h reaction time frame. Therefore, an agent is necessary to transport tin to the cold end reaction site. Furthermore, the temperatures reached during the reaction are suitable for the disassociation of tin iodide into tin and iodine.14 We find that placing Sn and P (without SnI4) at the cold end of the ampule and running the reaction under a normal b-BP growth temperature profile forms VP needles and ribbons (Supplementary Figure 8), not b-BP. Thus, iodine functions to transport Sn to the cold end reaction site and plays a role in the formation of BP. Since running the reaction with only red phosphorus leads to no transformation, it seems likely that phosphorus precipitation from a Sn-P melt upon cooling is an essential part of the formation process. Nucleation centers can reduce the kinetic barrier to this condensation. In contrast, the substrate does not appear to play a critical role in the growth process. Li et al. has suggested that clathrate may serve as nucleation centers for the growth of VP, which then converts to b-BP, 7095
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Figure 6. (a) Cross-sectional TEM HAADF image of BP along the armchair crystallographic axis showing SnIx crystal inclusions. (b) SnIx inclusions at higher magnification.
inclusions in the same plane of BP suggest that the inclusions solidify first and serve as a template for the lateral growth of BP. As the BP crystal extends beyond the lateral boundaries of the melt and makes contact with the substrate, the combination of the seed and the BP crystal begins to move with respect to the melt. Indeed, we often see amorphous residue on the surface of the wafer that terminates at the BP crystal (see Supporting Information). In this model, the width of the crystal is set by the diameter of the melt. This would mean that the melts in this work are much smaller than in the bulk growth recipe since, although the aspect ratio is similar, the BP film crystals are narrower than the b-BP crystals. The problem with this suggested mechanism is that under bulk growth conditions, the growth interval for b-BP is 400− 500 °C.13 However, stannous iodide has a melting temperature of 320 °C, which suggests that the thin-film growth process may be substantially different than that of the bulk growth process in spite of the similarity of the precursors. Multiple alternative growth mechanisms are known to exist for BP, and these change the growth temperature. These include metal with SnI4, bismuth, and metal alloys with I2.20,49 It has been shown that BP structures such as microbelts can be grown from Bi at temperatures as low as 270 °C.49 To better understand these results, consider the Sn:I:P ternary phase diagram (Supplementary Figure 10). We choose to use the diagram of Pfister et al. because it includes the compounds SnIP and Sn24P19.3I8.50 Next, we adopt the idea that the clathrate is unlikely to form due to the high I:P ratio in our experiments. SnIP dissociates at 600 K (323 °C), about the same temperature as the SnI2 solidification temperature. A possibility, therefore, is that this dissociation is responsible for the growth. In this model, most of the material from SnIP dissociation goes into the gas phase as SnI4 and P4. However, the cooling surface forms small SnI2 crystals which seed the growth of a few layers of BP. As the melt cools from the top, additional SnI2 nanocrystals are formed which nucleate additional layers of BP. The fact that iodine-deficient Sn24P19.3I8 dissociates at 740 K (467 °C), a temperature that is consistent with the bulk growth formation temperature suggests that the difference between the bulk and the thin-film growth methods is the identity of the intermediate growth compound. Additional iodine favors SnIP rather than Sn24P19.3I8.
however the evidence given for this was not conclusive. To test this hypothesis, we took large monoclinic VP samples grown using alkali metal rubidium in a sealed ampule with a max temperature of 550 and 50 °C thermal gradient.42 The monoclinic VP samples were exfoliated on to a silicon wafer and sealed in a quartz ampule with elemental iodine. Annealing exfoliated VP on a silicon wafer with iodine did not cause it to undergo a phase change to BP after 1 h at a temperature range of 450−540 °C. In contrast, Nilges’s work with in situ neutron diffraction that suggests that no intermediate exists;13 that is, bBP is formed directly from RP, although the mechanism is uncertain. This report received some confirmation from Li et al., who showed that the crystallized clathrate can serve as a nucleation site for the direct growth of VP.43 Our growth parameters for thin-film BP crystal growth (Figure 1) shows iodine plays a major role in forming thin films of BP. Increasing the magnification of the TEM image (Figures 4b and 6b) shows that the inclusions which always form parallelogram shapes are well incorporated into the existing BP lattice. The stacking directions of the BP planes are always aligned with the edge of the parallelogram, meaning that the inclusions are parallel to the stacking direction. EDS mapping shows that the crystals are formed with tin and iodine, and possibly phosphorus. Unfortunately, since the crystal inclusion dimensions are small and they are embedded in a phosphorus matrix, it is difficult to extract precise compositional data. However, the iodine concentration appears to be larger than the tin concentration. The data for four inclusions place the Sn to I ratio in the 1.1−1.4 range. EDS quantification is also difficult because metal halides are highly susceptible to radiation damage from the electron beam.44−46 Rapid damage was visually observed after collecting EDS maps, and halide desorption commonly occurs under increasing ebeam dosage, so the true iodine content is likely to be higher than measured. Thus, the measured composition of the inclusions is inconsistent with the clathrate, the proposed nucleation center for b-BP, but may be consistent with tin(II) iodide (stannous iodide or SnI2). Tin(II) iodide is a 2D material with c = 1.072 nm compared to the BP crystal with c = 1.047 nm, a mismatch of only 2%.47,48 This would explain the observation that the averages of the atomic positions in HAADF imaging are not disturbed to a significant degree by observing the interface between the inclusions and the greater BP lattice structure. In this case, the free surface of the melt will cool first due to the high thermal mass and excellent thermal conductivity of the silicon. The lattice registration of the BP and inclusions strongly suggests epitaxy. Furthermore, the relative size of the BP and inclusions as well as the fact that we never see two
CONCLUSION Direct growth of BP onto silicon and silicon dioxide substrate is achieved under mild growth (
