Open access peer-reviewed chapter

Defect Engineered 2D Materials for Energy Applications

By Sai Sunil Kumar Mallineni, Sriparna Bhattacharya, Fengjiao Liu, Pooja Puneet, Apparao Rao, Anurag Srivastava and Ramakrishna Podila

Submitted: November 17th 2015Reviewed: June 14th 2016Published: August 31st 2016

DOI: 10.5772/64605

Downloaded: 1088

Abstract

Two-dimensional (2D) materials display unique properties that could be useful for many applications ranging from electronics and optoelectronics to catalysis and energy storage. Entropically necessary defects are inevitably present in 2D materials in the form of vacancies and grain boundaries. Additional defects, such as dopants, may be intentionally introduced to tune the electronic structure of 2D materials. While defects are often perceived as performance limiters, the presence of defects and dopants in 2D materials results in new electronic states to endow unique functionalities that are otherwise not possible in the bulk. In this chapter, we review defect-induced phenomena in 2D materials with some examples demonstrating the relevance of defects in electronic and energy applications. In particular, we present how the (i) N-dopant configuration in graphene changes the electron-phonon interactions, (ii) zigzag defects and edges in graphene increase the quantum capacitance to improve energy density of graphene-based supercapacitors, and (iii) charged grain boundaries in exfoliated Bi2Te3 preferentially scatter low-energy electrons and holes to enhance the thermoelectric performance.

Keywords

  • Engineering defects
  • 2D materials
  • Quantum capacitance
  • Energy storage

1. Introduction

Two-dimensional (2D) materials have intrigued physicists and material scientists for many decades due to an abundance of unusual physical phenomena that result from the confinement of charge, heat, and entropy flow to a plane [1]. For example, the ingenious harness of quantum mechanical phenomena, particular to lower dimensionality in graphene, has resulted in intriguing observations such as the quantum Hall effect at room temperatures, quantized optical transmittance, nonlocal hot carrier transport, and Klein tunneling [2, 3]. Despite such fundamental breakthroughs, the potential of 2D materials has not yet completely manifested into practical devices due to material limitations [1, 4]. For instance, the lack of a band gap resulted in serious limitations for using graphene in electronics [4]. Defects in material science and engineering are often perceived as performance limiters, but in the case of 2D materials, defect engineering could provide a way to overcome many roadblocks and forge new frontiers. In this regard, others and we have shown that defects in 2D materials (e.g., dopants, vacancies) can provide an excellent handle to control material properties [58]. Specifically, we have shown that defects such as vacancies and N dopants in graphene could be used to control the electron-electron and electron-phonon scattering pathways [8]. These results provided critical breakthroughs for improving the quantum capacitance of graphene and doping graphene without compromising its intrinsic characteristics [6]. Defects also play a vital role in improving the properties of the so-called “beyond graphene” 2D materials. Previously, we used spark plasma sintering (SPS) to introduce charged grain boundaries (GB) in 2D Bi2Te3 for improving its thermoelectric (TE) figure of merit and compatibility factor [5]. Similarly, it has been demonstrated that defect engineering in 2D materials could improve many qualities ranging from electronic levels, conductivity, magnetism, and optics to structural mobility of dislocations and catalytic activities [9, 10]. As discussed in this chapter, defect engineering in 2D materials leads to the discovery of potentially exotic properties, which can enable unprecedented technological applications. In particular, we present how dopants and defects in (i) graphene could be used for optical and electrochemical energy storage applications and (ii) 2D Bi2Te3 could be controlled for enhancing its thermoelectric efficiency.

2. Nitrogen dopants for tuning the electronic and optical properties of graphene

Graphene is an ideal platform for many optoelectronic devices due to its distinctive combination of high electron mobility (μ), optical transparency, and gate/dopant-tunable carrier density [2]. However, to truly harness the potential of this combination and make graphene-based efficient optoelectronic devices a reality, optical and electronic properties of graphene must be tuned via substitutional doping [8]. While doping graphene with boron (B) or nitrogen (N) can tune the Fermi energy (EF) and lead to p- or n-type graphene, it also compromises the inherently high electron mobility in doped graphene, which is useful for electronic applications [1113]. Furthermore, the recombination rate of photogenerated carriers is also known to decrease with the presence of defects and dopants limiting applications of graphene in optoelectronics [8]. Indeed, the focus of many synthesis efforts in graphene has been to achieve large-area “defect-free” graphene for electronic applications because carrier-defect scattering limits the electron mobility—an important parameter for high-speed electronics [1417]. Contrary to this conventional wisdom, we recently demonstrated that the defect configuration (in particular, N-dopant configuration in bilayer graphene lattice shown in Figure 1a) is more important than the defect concentration for increasing carrier concentration without compromising photogenerated carrier lifetime [8].

Figure 1.

(a) Possible nitrogen dopant configurations in graphene lattice include pyridinic (green), pyrrolic (purple), and graphitic (red). While graphitic dopants (red) are purely substitutional, non-graphitic dopants (pyridinic and pyrrolic) result in additional defects such as vacancies or pentagons. (b) A schematic representation showing the linear energy dispersion for single-layer graphene and the arrow indicates excitation of electrons from the valence to the conduction band. (c) Under intense photoexcitation, the nonequilibrium carrier distribution (in the E-k space depicted by the energy dispersion shown in (b)) results in an initial rise in the transmitted intensity. The carriers (electron and hole distribution shown in blue and orange, respectively) equilibrate by carrier-carrier scattering on a timescale τ1. Subsequently, the carrier thermalization and decay occur through carrier-phonon scattering on a timescale τ2.

The electronic and optical properties of a single-layer graphene (SLG) can be described in terms of massless Dirac fermions with linear dispersion near the Fermi energy (Figure 1b). The semimetallic nature and electronic band structure of SLG allow for the photogeneration of electron-hole pairs at any wavelength in the visible-light spectrum [3]. This property is critical for many wide-bandwidth optoelectronic applications. As shown in Figure 1b, incident light excites electrons from the valence band (orange) into the conduction band (purple). Shortly after photoexcitation, incident photon-electron interactions create an out-of-equilibrium electron distribution (purple in Figure 1c), which initially relaxes on an ultrafast timescale (τ1 ~ 100–300 fs) to a hot Fermi-Dirac distribution and subsequently cools via phonon emission or defect scattering (τ2 ~ 1–2 ps) in graphene [8, 18]. In optoelectronic devices, when photoexcited electrons are scattered by phonons or defects, energy transferred to the lattice is dissipated as heat decreasing the net energy transported through charge carriers to drive a circuit. In the current scenario of graphene optoelectronic devices, a critical challenge is to increase the net charge carrier density and quench electron-defect relaxation pathways to extend photogenerated carrier lifetime.

The influence of defects on photogenerated carriers could be accounted by the inclusion of an extra term (A) in the expression for carrier scattering rate τ−1 (N) = A + BN + CN2, where A represents nonradiative recombination, usually due to defects or traps, B represents radiative recombination, N is the carrier density, and C represents Auger recombination. The photogenerated carriers are quickly cooled to ground state through scattering by defects (represented by A) in addition to the existing carrier-carrier and carrier-phonon scattering (term B) and Auger recombination (term C). In the context of optoelectronic applications, it is imperative to identify ideal dopant concentration and configuration in graphene for which A in the carrier scattering rate equation is minimized. As mentioned earlier, heteroatomic doping can tune EF and alter density of states (DOS) and thereby modify its electronic and optical properties. For instance, N-doping in chemical vapor deposition (CVD) graphene resulted in a bandgap due to the suppression of electronic density of states near the Fermi level and a consequent reduction in μ [19, 20]. As mentioned earlier, it is well known that N atoms can be substitutionally doped in the graphene lattice either in the pyridinic, pyrrolic, or graphitic configurations (see Figure 1a). Our density functional theory (DFT) calculations previously showed that all the three dopant configurations (viz., pyridinic, pyrrolic, and graphitic) are stable structures with a positive energy (>9.5 eV) released during the formation with graphitic dopants exhibiting the highest stability (Figure 2) [8].

Figure 2.

The lattice structure (top three panels) of graphitic (a), pyridinic (b), and pyrrolic (c) defects in graphene along with their electron density, obtained from our DFT calculations, shown in the bottom panels. The energy released on the formation of structure from free atoms for all the configurations was found to be positive (graphitic, 10.22 eV; pyridinic, 9.77 eV; and pyrrolic, 9.55 eV) confirming the stability of N-doped configurations.

2.1. CVD synthesis of N-graphene

Previously, we employed atmospheric pressure chemical vapor deposition (CVD) method for growing N-graphene [7]. This CVD set up consisted of a Cu foil loaded inside a 1 in. quartz tube at a temperature of 1,000 °C. Methane gas was used as the carbon source for graphene growth, while acetonitrile (AN) and benzylamine (BA) were used as precursors for N dopants in varying concentrations. The reaction was carried out under inert atmosphere by passing a mixture of Ar and H2 through the quartz tube reaction chamber. In particular, 450 sccm of Ar and 50 sccm of H2 were used, and 2 sccm of methane was bubbled through the mixture of BA and AN. The volume percent of BA and AN varied in the ratio of 0:1, 1:1, and 3:1. Accordingly, the obtained samples were labeled S1, S2, and S3, respectively. Interestingly, we found that the N-dopant configuration (viz., graphitic, pyrrolic, and pyridinic) could be controlled using the ratio of BA to AN precursors.

2.2. The effects of N dopants in graphene: X-ray and Raman spectroscopy

We observed a strong correlation between the N-dopant configuration and the accompanying vibrational properties of N-doped CVD graphene: the N atoms bonded in the non-graphitic configurations (pyridinic and pyrrolic, observed using X-ray photoelectron spectroscopy or XPS) resulted in intense Raman disorder bands unlike the N atoms bonded in the graphitic configuration, even though the concentration of N dopants was higher in the latter case [7].

As shown in Figure 3a, we identified XPS peaks corresponding to graphitic, pyrrolic, and pyridinic configurations [21]. For pyridinic configuration, the N1s peak positions reported in the literature are usually in the range 398.1–399.3 eV. Similarly, pyrrolic configuration gives rise to peaks in the range 399.8–401.2 eV, while the peak around 400.5 eV (blue colored) is associated with the graphitic configuration. The orange-colored peak at ~401.5 and 406 eV may be attributed to different nitrogenated adsorbents [20, 21]. XPS results confirmed that the atomic percentages of nitrogen in S1, S2, and S3 were 0.2, 2.5, and 3.8 %, respectively [7]. It is important to note that S1 and S3 showed more non-graphitic N dopants compared to S2, which was purely graphitic doping. In order to further understand the effect of various nitrogen-doping configurations and concentrations on the electronic structure of graphene, we performed Raman spectroscopy of the samples S1, S2, and S3. The Raman spectrum of graphene displays four important bands [22, 23]: (i) the disorder or D band appears ~1350 cm−1 due to the presence of defects such as edges, grain boundaries, or any other type of defects including dopants in the graphene lattice; (ii) in some studies, researchers have also reported the presence of D′-band ~1600–1625 cm−1 in the Raman spectrum of highly disordered graphene [22]; (iii) the graphitic G band ~ 1585 cm−1 arises due to doubly degenerate optical phonon modes at the Brillouin zone center. It is a first-order Raman scattering process, and (iv) the 2D band ~ 2700 cm−1 is a consequence of second-order Raman scattering process involving intervalley scattering of in-plane transverse optical (iTO) phonons. Unlike the D band, which requires the presence of a defect to conserve the momentum of scattered electron along with an iTO phonon, 2D band does not require the presence of any defect due to the involvement of two iTO phonons and is always present in Raman spectra of both pristine and doped graphene [22]. The shape and width of the 2D band are sensitive to the number of layers in graphene. For example, 2D band in single-layer graphene (SLG) can be fit into a single Lorentzian peak, while for bilayer graphene, 2D band can be deconvoluted into four sub-peaks [23].

As seen in Figure 3b, the Raman spectra of pristine graphene samples did not exhibit strong D band in our studies. While samples S1 and S3, which contain non-graphitic doping configuration of nitrogen, showed strong D bands, the D band in sample S2 (graphitic) is similar to that in pristine sample despite higher dopant concentration (~2.5 %). These results are consistent with our observations in the XPS spectra shown in Figure 3a. When nitrogen atoms enter the graphene lattice in non-graphitic configuration, vacancies are needed and result in armchair-type edges. Previous reports showed that armchair edges in graphene allow intervalley scattering of iTO phonons in the Brillouin zone unlike zigzag edges [22] and thus increase the intensity of the D band as in samples S1 and S3.

Figure 3.

(a) XPS spectra for N1s line. The colored peaks represent the deconvolution of N1s peaks. The resolution of the spectrometer was 0.5 eV. The peak at 400.5 eV corresponds to graphitic configuration. Peaks at 398 and 400 eV correspond to non-graphitic. Raman spectra of pristine and N-graphene are shown in (b) and (c). As seen in (b), the D and D′ bands increase in intensity for non-graphitic samples S1 and S3. The deconvolution of 2D band in (c) suggests that graphene samples are bilayer.

From the line-shape analysis of Raman 2D band (Figure 3c), we confirmed that our CVD-grown graphene samples are predominantly bilayers. As seen in Figure 3c, maximum downshift in 2D band (25 cm−1) was observed for sample S3 with relatively large dopant percentage (~3.5 %). On the other hand, sample S2 (graphitic configuration) showed little downshift in 2D band compared to sample S1 in spite of having higher dopant concentration. 2D band in S1 showed a downshift of ~10–15 cm−1 even in the presence of low dopant concentrations (~0.2 %). These differences in the 2D band shift in the Raman spectra can also be attributed to the nature of the dopant environment. For example, in samples S1 and S3 that are non-graphitic in nature, due to lattice symmetry breaking, electronic structure of graphene is strongly perturbed leading to possible renormalization of electron and phonon energies. Such a renormalization in electron energies results in a concomitant downshift in phonon energies of 2D band [7].

2.3. Nonlinear optical studies of N-graphene

We further explored the influence of defects on the carrier scattering rate using pump-probe (PP) spectroscopy [8]. The differential transmittance (ΔT/T) was obtained by taking the ratio of pump-induced change in the probe transmittance (ΔT) at a time t after the pump excitation to the probe transmittance (T) in the absence of a pump (Figure 4a). The initial response is an incident pulse-width-limited rise in the transmitted signal immediately after the zero delay (t = 0), which eventually decays in an exponential manner. The best fit to the PP data was obtained with a bi-exponentially decaying function, ∆T/T = A1 exp(−t/τ1) + A2 exp(−t/τ2) with two distinct time scales: a fast component (τ1) corresponding to the intraband carrier-carrier scattering and a slower component (τ2) corresponding to carrier-phonon scattering (discussed before in Figure 1b). Both τ1 and τ2 values obtained from the best-fitted curves for the samples are tabulated in Figure 4b. As shown in Figure 4b, the carrier-carrier and carrier-phonon relaxation times do not decrease monotonically with increasing N content, akin to the Raman features described in Figure 3. Clearly, the dopants present in the non-graphitic configurations exhibit much faster relaxation times relative to dopants present in S2 with graphitic bonding or pristine graphene. Indeed, carriers in sample S1 decay much faster (within fs), despite lower N content, due to the presence of extended defects that lead to increased contribution from carrier-defect scattering. More importantly, N dopants in graphitic configuration do not affect the carrier recombination times even at >1 % doping, while pyrrolic/pyridinic configurations drastically decrease carrier lifetimes at concentrations as low as ~0.2 %.

Figure 4.

(a) Time-resolved differential transmission spectra of pristine and N-doped graphene samples S1, S2, and S3, respectively, obtained through pump-probe spectroscopy. The inset shows parabolic energy dispersion for bilayer graphene, where the arrow indicates excitation of electrons from valence to the conduction band. The solid lines represent bi-exponential fits based on the above equation. (b) Tabulated values of relaxation times τ1 and τ2, corresponding to intraband carrier-carrier scattering and carrier-phonon scattering, respectively, obtained from the fits.

3. Defects in graphene for energy storage

The increasing global energy demands have spurred a rigorous search for new renewable energy sources. In recent times, fuel cells, photovoltaic devices/solar cells, and various other renewable energy sources have received much attention and are all promising candidates for clean energy production. However, today’s batteries and capacitors, which are the main components for energy storage, cannot meet the world’s demand for combined power and energy densities [2427]. As an example, the plot in Figure 5 shows the general performance metrics for commercially available formats of charge/energy storage devices. This plot depicts electrical capacitors (on the top left hand side) with a fast response time and high-power-density batteries and fuel cells (on the bottom right hand side), which exhibit a large energy density due to the chemical/ionic basis of their reactions. The drawback of electrical capacitors is their inability to store large amounts of energy, while the batteries are incapable of fast charge/discharge cycles due to the slow nature of the ion diffusion processes. This gap between capacitors and battery performance has been a major roadblock in electrochemical energy storage. Electrochemical double-layer capacitors (EDLCs) or supercapacitors have been proposed to bridge the gap between these disparate devices by incorporating elements of both technologies [12, 2833]. The charge is stored in the electric double layer, which forms at the electrode/electrolyte interfaces and leads to a double-layer capacitance (Cdl). The specific energy (E) of the EDLC may be expressed in terms of the total measured device capacitance (Cmeas) and the operating voltage (V) as E = (1/2)CmeasV2. Unlike traditional Li-ion batteries, EDLCs can be reliably used for hundreds of thousands of cycles since their charge storage mechanism does not involve ion motion and the consequent chemical irreversibility. Although the EDLCs are superior (say, to batteries) in terms of long-term cyclability, they suffer from poor energy density in that energy (or charge) is only stored at the surface/interface rather than within the bulk of the material.

Figure 5.

(a) Ragone plot showing specific power versus specific energy for common electrical energy storage devices. Supercapacitors are expected to bridge the gap between batteries and capacitors and impact nearly every area of electrical energy usage. For practical applications, the energy and power densities indicated by the star are needed. (b) Cyclic voltammetry (CV) characterization of plasma processed FLG samples (in 0.25 M TBAHFP dissolved in a 1 M acetonitrile). The area enclosed by the CV curves was used to parametrize the Cmeas which increases with plasma power (indicated on the figure). (c) A close to threefold enhancement in the Cmeas (left axis) and the contributions of the computed Cq (right axis), i.e., 1/Cmeas = 1/Cdl + 1/Cq, as a function of the plasma power.

Nanocarbons including carbon nanotubes and graphene have been widely used as an electrode in EDLCs due to their high surface area (~2000 m2/g), modest electrical conductivity, electrochemical stability, and open porosity [25, 27]. However, the performance of the carbon-based EDLCs (particularly, graphene) is fundamentally limited by the so-called quantum capacitance (Cq), which is defined as Cq = e2DOS(EF) with e being the charge of an electron [6]. An intrinsically small DOS (EF) in graphene results in a small serial Cq, which diminishes the total device capacitance value (1/Cmeas = 1/Cdl + 1/Cq) in EDLCs [6, 28]. Although graphene has high surface area and consequently high Cdl, the total EDLC energy is limited by small Cq. Defects can be advantageous for alleviating the limitation of Cq by increasing the DOS (discussed later in Figures 5b and c). In our previous work [6], we synthesized few-layer graphene (FLG) on Ni foil substrates through chemical vapor deposition and subsequently subjected to argon-based plasma processing to intentionally induce charged defects (see the inset in Figure 6). Argon was chosen because the constituent ionic species are limited to Ar+, implying relatively simple plasma chemistry.

Figure 6.

Enhanced plasma processing applied to the pristine sample results in a substantial intensity enhancement of the D- and D′-peaks as seen in the Raman spectra (normalized to the G-peak). An increase in power, say to 50 W, may result in irreversible changes due to graphene removal. The inset schematic shows the difference between armchair and zigzag defects (created using plasma processing) in the graphene lattice.

As shown in Figure 6, we found that the D band in the Raman spectrum of graphene increased with increasing power of plasma etching due to the introduction of new structural defects such as pores, which contain both armchair and zigzag edges. An important attribute of zigzag defects is that they may be electrically active and could contribute to an enhanced DOS much more than the armchair-type edge defects (which contribute less due to the two constituent carbon atoms belonging to different sublattices). We observed that the increase in plasma power resulted in a high device capacitance due to higher Cq arising from defect-induced DOS (EF). Indeed, we used cyclic voltammetry to quantify the changes in Cq and Cmeas (Figure 5b). The more than doubling of the Cmeas from 1.9 μF/cm2 (for the pristine sample) to 4.7 μF/cm2 (for the sample subject to 20 W plasma) is remarkable and suggests a novel means of substantially enhancing capacitance through defects. However, at higher plasma power (>20 W), high defect concentration results in poor electrical conductivity leading to a drop in Cmeas suggesting the importance of defect concentration in determining 2D material properties.

Thus, as evidenced by our data in graphene, the presence of defects does not necessarily deteriorate the material performance. Though there is only one way for a given material to be defect-free, there are many possibilities for it to be imperfect. While defect configuration is important in determining the mobility through carrier scattering rate, controlling defect concentration is critical for electrochemical applications. Accordingly, future efforts must be focused on finding new approaches to identify and control the right defect configurations (e.g., N in graphitic configuration to increase carrier concentration without compromising carrier scattering rates or mobility) and concentrations, which could improve material properties instead of dismissing all defects as detrimental for carrier mobility.

4. Defects in 2D bulk materials for thermoelectric power generation

Thermoelectric (TE) materials have the potential to reduce global energy crisis and global warming effects by converting waste heat to electricity. As of 2005, the world energy usage was ~15 terawatts of energy, of which ~90 % was first converted to heat and the remainder ~10 % of energy was utilized [24]. In general, power plants and the transportation industry are the two main sources of heat energy losses that contribute to global warming. In recent years, prototype car models developed by automobile industries BMW and Ford have successfully transformed the waste heat from car exhausts to electricity using thermoelectric power generators, thus improving the fuel efficiency [24, 34].

A basic thermoelectric energy conversion module consists of n and p-type semiconducting materials, connected electrically in series and thermally in parallel [35]. The maximum thermoelectric efficiency is a product of the Carnot Efficiency and a term consisting of ZT or the thermoelectric figure of merit, which embodies interrelated material parameters, as given by ZT = α2σT/k, where α is Seebeck coefficient or thermopower, σ is the electrical conductivity, and k(= kE + kL) is the total thermal conductivity comprised of electronic (kE) and lattice contributions (kL), respectively. The main challenge of improving the energy conversion efficiency and consequently the ZT of thermoelectric materials is the inherent coupling between the electrical conductivity and the Seebeck coefficient. In recent years, there has been significant scientific progress in the field of thermoelectrics with the use of nanostructured materials (e.g., superlattices, nanowires, and nanocomposites) that have simultaneously increased the power factor (the numerator of ZT) and reduced the thermal conductivity to achieve a high ZT > 1 [36].

4.1. Quantum confinement effects in 2D thermoelectric materials

In the early 1990s, Hicks et al. [37] predicted intriguing changes in transport properties upon lowering the dimensionality of existing bulk materials (e.g., from 3D to 2D) that were not observed in the corresponding bulk materials. A dramatic increase in the density of states (DOS) of low-dimensional materials was predicted that could increase the Seebeck coefficient and potentially decouple the electronic transport properties. Moreover, the presence of numerous interfaces in low-dimensional materials also increased phonon scattering effects that reduced the lattice thermal conductivity, thus introducing opportunities to independently vary all the parameters constituting the ZT [38]. In the low-dimensional thermoelectric materials, dramatic changes in the density of states were observed in the PbTe-based thermoelectric materials owing to the quantum confinement effects that led to further decoupling of the TE transport properties and an enhancement of ZT [39].

Nevertheless, it is challenging to fabricate low-dimensional materials for commercial thermoelectric applications and devices, which must be a cost-effective and facile process. In addition, the nanostructured thermoelectric materials have to be thermodynamically stable to retain the desired 2D properties over time, to make the devices reliable and long lasting. To be able to make use of the advantages of low-dimensional materials as well as robustness of the bulk materials, bulk nanomaterials or nanocomposites have been used to enhance the thermoelectric performance of existing thermoelectric materials such as SiGe and PbTe [40, 41].

Controlling the multi-scale microstructures via defect engineering and consequently the length scales of the electrical and thermal transport is essential for enhancing TE performance. However, material properties (thermopower, electrical, and thermal conductivity) dictating the ultimate compatibility factor and ZT are inherently coupled at the nanoscale. Thus, any efforts to improve only one particular property (e.g., thermopower) using microstructure changes are often futile as they inadvertently deteriorate remaining properties (e.g., electrical conductivity). In this section, we describe how charged defects in 2D materials can introduce low-energy carrier filtering and selective charge carrier scattering (e.g., hole excessively scattered compared to electrons) to improve ZT and compatibility factor in Bi2Te3.

4.2. Impact of charged grain boundaries in few-layered Bi2Te3

Discovered in the early 1950s by Goldsmid [42], Bi2Te3 is one of the most used and commercialized TE materials for room-temperature power generation and refrigeration applications [43]. The first TE refrigerator was designed using the p-type Bi2Te3, which was estimated to have a figure of merit ~ 0.48. In 2001, the thin-film superlattices were reported to exhibit enhanced ZT ~ 2.4 and ~1.4 at 300 K in the p-type Bi2Te3/Sb2Te3 and the n-type Bi2Te3/Bi2Te2.83Se0.17, respectively [44]. While a high ZT ~ 1.5 was realized in the bulk p-type Bi2Te3 via nanostructuring techniques such as ball-milling [45], melt-spinning [46, 47], and hydrothermal synthesis [48], it was found that these methods were ineffective for the n-type Bi2Te3. Like graphite, n-type Bi2Te3 is easily cleavable, and these traditional nanostructuring methods were too harsh and resulted in the deterioration of its basal plane (or in-plane) properties (Figure 7a).

Recently, Puneet et al. [5] utilized a novel technique of chemical exfoliation followed by spark plasma sintering (CE-SPS) in the n-type Bi2Te2.7Se0.3 that significantly improved the TE compatibility factor and stabilized the ZT peak at higher temperatures (300–450 K) as compared to the commercial ingot. Based on the studies of emerging two-dimensional (2D) materials (e.g., graphene), it may be expected that the electronic, thermal, and optical properties of chemically exfoliated n-type Bi2Te3 should exhibit properties that are different from the bulk. In the following paragraphs, we discuss the effect of the chemical exfoliation and spark plasma sintering of n-type Bi2Te3 as compared to the bulk commercial ingot.

The bulk bismuth telluride (Bi2Te3) exhibits a rhombohedral crystal structure belonging to the space group R 3¯m(D5), which is more commonly represented by a hexagonal crystal structure as shown in Figure 7a. The hexagonal unit cell of Bi2Te3 is composed of three quintuples with lattice constants a = 0.4384 Å and c = 3.036 Å, respectively. As shown in the figure, each quintuple consists of five atoms stacked in the order Te1-Bi-Te2-Bi-Te1 along the c-axis, which are bonded by ionic-covalent bonds. The Te1-Te1 layers between the two quintuples are held together by the weak van der Waals forces that make the Bi2Te3 easily cleavable. The n-type Bi2Te3 was obtained by partially doping Se at the Te1 or Te2 sites, which represent the Te atoms with two types of bonding.

Figure 7.

(a) A schematic representation showing Bi2Te3 rhombohedral crystal structure belonging to the space group R3¯ m(D5). Upon chemical exfoliation, the crystal breaks at the van der Waals gap. (b) A scanning electron micrograph showing the layered structure of bulk Bi2Te3 ingot. (c) A transmission electron micrograph showing chemically exfoliated quasi-2D Bi2Te3, which is repacked using spark plasma sintering (d) to achieve better thermoelectric performance through charged grain boundaries.

The nanostructuring of the commercial n-type Bi2Te3 ingot (see Figure 7b, purchased from the Marlow Industries, USA) was achieved by the chemical exfoliation method, using N-methyl-2-pyrrolidone (NMP) solution which was ultrasonicated using a tip sonicator (Branson 250) for 3, 5, 8, or 13 h, respectively. After centrifuging the solution and vacuum filtering the supernatant through a 0.45 mm nylon filter paper, the filtered powder was washed several times using deionized water to remove any residual NMP. The exfoliated nanolayers (see Figure 7c) were ~50 nm in thickness, which were then compacted using the spark plasma sintering (SPS, Dr. Sinter LabH-515S) technique at a holding temperature of 500 °C and an applied pressure of 30 MPa for 5 min under a dynamic vacuum. For the SPS process, the samples were loaded into graphite dies, and the pressure was applied using graphite rods. The resulting SPS-processed pellets were 12.5 mm in diameter and 2–3 mm in thickness, yielding samples with 98–99 % of the theoretical density (see Figure 7c).

The rapid densification technique by SPS has distinct advantages over other types of sintering techniques, such as the hot pressing. The SPS process is capable of sintering material powders within a very short time, in the order of minutes [49, 50]. As a result, it is possible to retain the metastable micro-/nanostructures of materials by limiting their grain growth and excessive diffusion during the sintering process. Furthermore, unlike other sintering techniques like hot pressing which uses furnace heating, only the graphite cylinder, rods, and sample are heated by the joule heating produced by pulsed dc in the SPS, which leads to even shorter processing times.

Figure 8.

(a) Charged defects near the grain boundaries (GB) can induce excess majority carriers (b) and preferentially scatter holes over electrons and low kinetic energy carriers (red arrow in panels a and c) due to engineered charged GB potential barriers shown in (c).

Contrary to the general understanding that defects in a crystal lattice are detrimental to the transport properties of materials, the defects in 2D materials are extremely useful and could be manipulated to generate controlled defects for novel and innovative applications. As observed in the few-layered bulk Bi2Te3, the localized positive charges in the grain boundaries introduced extra electrons in the material, thereby increasing the carrier concentration n (see Figure 3a in Ref. [5]). This increase in n was attributed to the injection of donor-like defects (Figure 8a and b), arising from positively charged anti-sites/Te vacancies on the exfoliated grain boundaries, due to chemical/mechanical exfoliation [51]. Moreover, these positively charged or interfacial charged defects acted as a potential barrier (see Figure 8c) to selectively filter out low-energy holes (or minority carriers) as shown schematically by Puneet et al. (see Figure 3 in Ref. [5]). These charged defects in the few-layered n-type Bi2Te3 thus shifted the onset of the bipolar (or two carrier) effects and consequently the maximum ZT value to higher temperatures, thus optimizing the ZT over a broader range of temperature. In addition, the thermoelectric compatibility factor of the few-layered n-type Bi2Te3 was significantly improved, which determined the ability of these materials to be segmented to other thermoelectric materials such as PbTe at higher temperatures, for operation over a broader range of temperature.

In summary, the CE-SPS processing of 2D Bi2Te3 leads to preferential scattering of electrons at charged grain boundaries and optimizes the band filling, thereby increasing the electrical conductivity despite the presence of numerous grain boundaries, and mitigates the bipolar effect via band occupancy optimization leading to an upshift in ZT peak (by ~100 K) and stabilization of the ZT peak over a broad temperature range of ~150 K. These changes in electrical and thermal transport led in turn to a more device-design friendly compatibility factor.

5. Conclusions

As exemplified by graphene and Bi2Te3, the presence of defects and dopants imparts the host material with new micro/quantum states or energy configurations that can strongly influence optical, electronic, and thermal properties. In addition to the above properties, S and F dopants are being explored to make graphene magnetic [10, 5255], while N dopants are expected to provide much higher enhancements in quantum capacitance without compromising electrical conductivity [28]. Similar to Bi2Te3, defects in other layered systems such as SnSe and TaSe2 could be engineered to achieve better thermoelectric performance [5662]. Although this chapter presented only some examples of defects in 2D materials, the same concepts also hold true for other 2D materials such as MoS2, WS2, and BN. Indeed, some properties (e.g., luminescence and catalytic activity) of these materials can be tuned using defects [6368]. As the design and development of new 2D materials are costly, complex, and limited due to the relatively poor air stability of many materials (e.g., silicene and phosphorene), the realization of desired properties and functionalities through control of defects (e.g., vacancies, dopants) in 2D materials is necessary. Though there is only one way for a given material to be defect-free, there are many possibilities for materials to be imperfect. The global scientific endeavors on understanding defects, such as the efforts presented in this chapter, provide a glimpse of the enormous potential of defects warranting further interdisciplinary research efforts.

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Sai Sunil Kumar Mallineni, Sriparna Bhattacharya, Fengjiao Liu, Pooja Puneet, Apparao Rao, Anurag Srivastava and Ramakrishna Podila (August 31st 2016). Defect Engineered 2D Materials for Energy Applications, Two-dimensional Materials - Synthesis, Characterization and Potential Applications, Pramoda Kumar Nayak, IntechOpen, DOI: 10.5772/64605. Available from:

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