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[MDR_Complete Document for Review_rev.doc](https://mdr.nims.go.jp/filesets/299f5387-edcc-4694-a9d7-9a82d32cb5ee/download)

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[Naoki Fukata](https://orcid.org/0000-0002-0986-8485), [Wipakorn Jevasuwan](https://orcid.org/0000-0001-9117-2497)

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[Creative Commons BY-NC-ND Attribution-NonCommercial-NoDerivs 4.0 International](https://creativecommons.org/licenses/by-nc-nd/4.0/)

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[Formation and characterization of Group IV semiconductor nanowires](https://mdr.nims.go.jp/datasets/ce227e92-6456-47ec-9b5e-f9760817e63b)

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Formation and characterization of Group IV semiconductor nanowiresAuthor Name1 and Author Name1,21 Department One, Institution One, City One, Country One2 Department Two, Institution Two, City Two, Country TwoE-mail: xxx@xxx.xx AbstractTo enable the application to next-generation devices of semiconductor nanowires (NWs), it is important to control their formation and tune their functionality by doping and the use of heterojunctions. In this paper, we introduce formation and the characterization methods of nanowires, focusing on our research results. We describe a top-down method of controlling the size and alignment of nanowires that shows advantages over bottom-up growth methods. The latter technique causes damage to the nanowire surfaces, requiring defect removal after the NW formation process. We show various methods of evaluating the bonding state and electrical activity of impurities in NWs. If an impurity is doped in a NW, mobility decreases due to the scattering that it causes. As a strategy for solving this problem, we describe research into core-shell nanowires, in which Si and Ge heterojunctions are formed in the diameter direction inside the NW. This structure can separate the impurity-doped region from the carrier transport region, promising as a channel for the new ultimate high-mobility transistor.Keywords: nanowire, doping, heterostructure1. IntroductionSemiconductor nanowires (NWs) are regarded as potentially important building blocks for various next-generation nanoscale devices such as transistors, sensors and solar cells [1-9]. One strategy for improving the performance of transistor devices is to reduce their size, and the use of NWs for this purpose is being researched for future transistor channels [1-3,10-16]. In addition, by forming a pn junction or heterojunction inside the NW, the NW itself can act as a nanoscale solar cell that has potential for use in highly efficient flexible solar cells, in which the recombination of photoinduced carriers is suppressed and light reflection from the NW surfaces is significantly reduced [5,9,17-31]. These effects increase short-circuit current density, resulting in higher efficiency. Furthermore, highly sensitive sensors that take advantage of the high surface area ratio of NWs are being actively studied [8,27,32,33].For transistor applications of semiconductor NWs, the most important elements are the control of NW formation, impurity doping for functionalization, and device formation processes. Many studies have reported on the control of the growth and formation of NWs [1,34-47]. The management of diameter and direction of NWs are important for device applications. Growth has been realized in almost all materials from Group IV semiconductors, centred on Si and Ge, to compound semiconductor materials. In the 2000s, bottom-up growth using metal catalysts, called the vapour-liquid-solid (VLS) growth mechanism [34,48,49] was the main growth method, but since then, reports of NW arrays formed by top-down methods with controllable size and position have increased.Impurity doping into NW structures is often performed during growth, but results from ion implantation, a technique that is commonly used in semiconductor processes, have also been reported [1,34,50-56]. In the latter case, electrical activation after the implantation of dopant atoms is an important challenge [52-56]. The electrical activity of the dopant atoms is often characterized using electrical measurements [10-16]. Various spectroscopic methods have also been applied for close evaluation of the dopant concentration, the bonding state and the electrical activity of impurities [57-65]. The distributions of atoms in NWs with high spatial resolution and high sensitivity can also be characterized by atom probe tomography (APT) [66-78].The problem with impurity doping is how to minimize impurity scattering inside the nano-scale NWs. At diameters of several tens of nanometers, the effect of impurity scattering becomes significant, and the mobility of the carrier decreases. To solve this problem, research attention has moved to core-shell type NWs in which a heterojunction of Si and Ge is formed in the diameter direction [78-98]. Core-shell NWs composed of Si and Ge have a Type II band structure, which allows the band offset to separate the impurity doping region from the carrier transport region. This structure is a high electron mobility transistor (HEMT)-type device which is being studied intensively for use in compound semiconductors [99]. It has an exceptional structure that can achieve high mobility. We call these devices NW-HEMTs. Axial heteronanowires with heterostructures along the growth axis have also been studied [100,101].,)1()()(220eew++=qIIIn this paper, we describe the research that has been conducted to apply Group IV semiconductor NWs to channels of next-generation vertical transistors. The following are reviewed: Control methods of growth and formation of Si, Ge, and Si-Ge core-shell NWs; control of impurity doping and its evaluation; and results of demonstration of hole gas accumulation in the carrier transport layer for realization of NW-HEMTs.2. Growth and formation of NWs2.1 Bottom-up and top-down methodsSemiconductor NWs have been grown and formed in various ways, and the size, shape, and arrangement of nanowires are controlled according to their purpose and application. Figure 1 shows scanning electron microscope (SEM) images of Si nanowires formed using three different methods and a schematic diagram of their growth and formation processes. Figure 1(a) shows the growth process of SiNWs by laser ablation, and Figure 1 (b) shows an SEM image. Research into SiNWs using laser ablation was the experimentation that triggered the NW research, as reported by the Harvard group in 1998 [102]. The metal catalyst was iron (Fe). The eutectic temperature of Fe-Si is about 1200 °C, and the growth of SiNWs requires a high temperature of about 1200 °C. In our group, Ni, whose eutectic temperature with Si is about 970 °C, was selected as the catalyst for the growth, and the alloy target of Si-Ni was laser-ablated using a frequency-doubled NdYAG laser (532 nm, 7 ns pulse width, 10 Hz, 100 - 200 mJ/pulse) in the temperature range of 1000 -1200 °C in flowing argon (Ar) gas at 50 sccm [103,104]. An Ar gas pressure of 500 Torr was the optimal under our experimental conditions. In this method, Si and Ni atoms evaporate atomically from the target (vapor state), after which Si and Ni atoms form nanoclusters under the pressure of the Ar gas [103,104]. Since this clustering takes place at a temperature of 1000 - 1200 °C, the nanoclusters become droplets, i.e., they are in a liquid state. Because the Si-Ni target contains a high proportion of Si, Si is preferentially incorporated into the nanocluster in the liquid state [103,104]. The nanocluster then becomes supersaturated with Si. When supersaturation occurs, Si precipitation takes place (solid state). Once precipitation occurs, precipitation continues in the precipitated portion, resulting in one-dimensional precipitation and NW growth. In the magnified part of the SEM image, the presence of a catalyst at the tip indicates that the nanowires have grown via VLS growth. The diameter of the NW can be controlled by the ratio of the metal catalyst in the target and the growth temperature: thinner NWs can be grown by lowering the ratio of the metal catalyst and lowering the growth temperature [104].Figures 1 (c) and (d) also show the results of bottom-up growth using a metal catalyst, on which SiNWs were grown by introducing a source gas via CVD. In this case, colloidal Au was applied to the Si substrate, and SiNWs were grown by setting the substrate temperature at 360 - 500 °C [91]. The orientation of the Si substrate used for growing the SiNWs is < 111 >, and the orientation of the SiNWs is also < 111 >, so they are vertically oriented. This is because growth in the < 111 > axis direction is most energetically advantageous when the diameter of the SiNWs is 20 nm or more [35]. The NW diameter can be controlled by the size and growth temperature of the colloidal particles used as metal catalysts. In short, thinner NWs can be grown by reducing the size of the colloidal particles and lowering the growth temperature.Figure 1 (e) shows a method for forming SiNW arrays using a method that combines top-down nanoimprint lithography and Bosch etching, and 1 (f) shows SEM observation results of the formed SiNWs. In the nanoimprint lithography method [97,105], a pattern is formed on a photoresist using a quartz mold to form a metal nanodot array. This metal nanodot array serves as a mask during Bosch etching, and a vertically aligned SiNW array can be formed on a Si substrate [97,105]. Bosch etching is a kind of reactive ion etching (RIE), and two gases, (SF6 and C4H8) are generally used. SF6 is used for Si etching and C4H8 is used to passivate the SiNW surface. The inductively coupled plasma (ICP) powers were 125 W for SF6 etching and 100 W for C4H8 passivation to obtain a smooth NW surface. The advantage of nanoimprint lithography is that rapid and repeated patterning can be applied over a relatively large area. On the other hand, as a disadvantage, it is difficult to reduce the diameter of the NWs to 50 nm or less using current techniques. Although the NW diameter can be reduced to 50 nm or less by using nanoimprint lithography rather than electron beam lithography, the process time is long and the simplicity of the nanoimprint lithography method is lost. It is necessary to use the former if formation over a large area is required and the latter when a fine pattern is needed.2.2 Surface defects and diameter controlConsidering future device applications, NWs should be formed using a top-down approach that is advantageous for NW alignment and size control. There are two problems to be solved in SiNW arrays formed on a Si substrate using the combined method of nanoimprint lithography and the Bosch etching method as shown in Figure 1(f). First, when Bosch etching, a RIE method, is used, defects due to damage appear on the surface of the SiNWs. This has a negative impact on all devices being considered for NW use. Figure 2 shows the results of an investigation of damage to the SiNW surface. As shown in the SEM images of Figures 2 (a) - (c), the surface to be annealed in an Ar atmosphere at 900 ˚C for 10 min was rough. However, this roughness was not observed when using an H2 gas atmosphere. The result of electron spin resonance (ESR) measurements taken to investigate damage to the SiNW surface by Bosch etching is shown in Figure 2 (d). The measurements were made at 4.2 K using an X-band ESR spectrometer with a magnetic field modulation of 100 kHz. The ESR signal with a g-value of 2.005 is assigned to dangling bond-type defects in the damaged layer on the SiNW surface caused by Bosch etching. The contribution of interfacial defects (Pb centers) formed at the interface between Si and the native oxide film is small [60,106]. The defect signals were observed for samples before annealing, meaning that the SiNW surfaces had been damaged by Bosch etching. Annealing in an Ar atmosphere reduces the signal intensity of defects, but the reduction rate is low. This low reduction rate is likely related to the surface roughness generated by Ar annealing. However, when annealing was performed in an H2 gas atmosphere, the ESR signal of this defect completely disappeared, and no surface roughness was observed. These results indicate that the surface shape of the SiNWs after Bosch etching and the removal of defects in the introduced damaged layer are closely dependent on the annealing atmosphere. Figure 2 (e) shows the different effects of the two ambient gases. White dots represent defects caused by damage. If there are many defects on the surface, they will aggregate during high temperature annealing. This is because the aggregated defects are energetically stabilized. Therefore, Ar annealing cannot effectively reduce the defects: instead, the surface becomes rough. On the other hand, when annealing in a hydrogen atmosphere, since hydrogen atoms are introduced during the annealing process, defects are ultimately removed from the NW surface because they inhibit the assembly of defects already present in the surface of the damaged layer. The differences seen with different annealing atmospheres can be explained by these mechanisms. The minimum size of a pattern formed by nanoimprint lithography is about 100 nm, and that with electron beam lithography is generally several tens of nanometers, which is larger than the intended size of transistor channels. Further, since a damaged layer is present on the surface of NWs formed by the top-down method, they cannot be used as it is, because this layer significantly degrades the device’s characteristics. Figure 3 shows the result of an experiment in which the NW diameters are reduced by removing the surface-damaged layer that had formed while using the top-down method. To remove the surface damage and reduce the diameter of the NWs, thermal oxidation was performed at 950 °C with O2 gas flow at 20 sccm/s, followed by etching in dilute HF solution. This oxidation etching process reduced the NW diameters to a minimum of about 50 nm [106]. Beyond that, the high aspect ratio of the NW structure may cause the NW to bend during processing in dilute HF solution during etching, and further diameter reduction would make it difficult to maintain upright NWs. The result of ESR measurements is shown in Figure 3 (c). When the oxidation treatment was carried out in an oxygen atmosphere, the ESR signal caused by the defect in the NW surface layer disappeared and was replaced by a new signal, observed at the position of g value of 2.000. This is an ESR signal caused by an oxygen deficiency defect in the oxide film and is called the E' center [107-111]. It indicates that the oxide film formed by the oxidation treatment is defective. After the oxide had been removed with dilute HF, the ESR signal indicating the oxygen-deficient defect disappeared and the ESR signal appeared again at the g value of 2.005. This means that even if the oxide film formed on the surface of the SiNW is removed, a new natural oxide film is formed in the atmosphere. Interface defects are inevitably formed at the interface between the native oxide film and the SiNW crystal. The observed ESR signal is due to interfacial defects called Pb centers [60,112,113]. Pb centers are well-known Si dangling bond-type defects (Si3 ≡ Si·) formed at the Si/SiO2 interface. The above results demonstrate the effectiveness of oxidation etching for removing the damaged layer and reducing the diameter of NWs.3. Impurity doping of NWs3.1 B doping in SiNWsThe nanometer-scale size of the material complicates impurity doping for carrier control, and evaluation of the bonding state and electrical activity of dopant impurities introduced into nanostructures is an important and challenging task. Electrical characterization is generally used to evaluate doping. However, electrical characterization requires electrode formation on the NWs, and the resulting properties are complex because they depend on the electrode formation process. Non-destructive evaluation, such as by using spectroscopic techniques, is therefore necessary. The results of Raman spectroscopy measurements are shown in Figure 4 (a) - (c). Micro-Raman scattering measurements were performed at room temperature (RT) using a 100x objective and a 532-nm excitation light. The excitation power was set at about 0.02 mW to prevent local heating effects from the excitation laser [103,104,114]. In Figure 4 (a), peaks are observed at about 618 cm–1 and 643 cm–1. Since these peak shifts coincide with the peak shifts of boron isotopes in Si, and the intensity ratio of the peaks also coincides with the natural abundance ratio of B isotopes (11B:10B = 4:1), the observed peaks can be respectively identified as localized vibrational peaks for 11B and 10B doped in SiNWs [50-52,59]. When hydrogen is introduced into p-type and n-type Si, a phenomenon called hydrogen passivation occurs [115-124]. Boron as an acceptor is made electrically inactive by hydrogen passivation. Because hydrogen enters the bond center (BC) site between Si and B atoms [115-121], it has been reported in bulk B-doped p-type Si that the local vibrational peak for B shifts slightly to the higher wavenumber side, that is, a slightly broader peak appears at the higher wavenumber side of the local vibrational peak for 10B [118]. As a result of introducing hydrogen atoms into B-doped SiNWs by hydrogen atom treatment (HAT) [124], a slightly broader peak was observed on the high-wavenumber side of the local vibrational peak for 10B, which is similar to that reported for bulk B-doped p-Si (Figure 4 (b)). This result supports the identification of the observed peaks at about 618 cm–1 and 640 cm–1 as local vibrational peaks for 11B and 10B. It is also important to investigate whether B introduced into SiNWs is electrically activated. The electrical activation of dopant impurities can also be evaluated by Raman spectroscopy [50-52,59]. The change in the Si optical phonon peak caused by B doping and the excitation wavelength dependence during Raman spectroscopy are shown in Figure 4. It can be seen that the Si optical phonon peak is asymmetrically broadened by B doping and also shows a downshift. These asymmetric broadening and downshift behaviors are further augmented by increasing the B doping concentration and by lengthening the excitation wavelength during Raman spectroscopy. This phenomenon can be explained by the Fano effect [4-52,59,125]. In this case, the Fano effect is caused by interference between the transition between discrete levels of phonons and electron transitions in the continuum between the heavy holes and the light holes in the valence band due to the high concentration of B doping. It is likely that asymmetric broadening of the phonon peak is generated as a result of this interference. The Fano effect means that the transition occurs in the valence band, which means that the B doped in the NWs was electrically activated. The electrically active B concentration can be estimated by analyzing the asymmetric broadening caused by the Fano effect. Fano's equation for fitting takes the following form. (1)where ω is the wavenumber, I0 the prefactor, q the asymmetry parameter, and ε is given by ε = (ω-ωp)/Γ, where ωp is the phonon wavenumber and Γ is the linewidth parameter. Here, the electrically active B concentration can be estimated from the q and Γ parameters obtained by fitting.Infrared absorption spectroscopy is a technique similar to Raman spectroscopy. It is difficult to detect very small amounts of dopant impurities in NWs using the infrared absorption spectrometers found in typical laboratories. To overcome this problem, results obtained by micro-FT-IR measurements using the IR synchrotron radiation (IR-SR) beamline at SPring-8 BL-43IR in Hyôgo, Japan were employed [126]. The IR-SR measurements were performed at 4.2 K. The brilliance of the IR-SR at SPring-8 is 104 times greater than that of ordinary black body light. In fact, the focal size of IR-SR is 10 times smaller and the light density is 102 times higher than that of black body light at the focal point of a conventional IR microscope. This means that the optical spectrum in a tiny region can be measured under high-contrast conditions. The IR-SR absorption peaks were observed at 4.2 K using a continuous-flow liquid-helium cryostat with KRS-5 windows. The spectral resolution was 2 cm–1. Figure 4 (d) shows the results of IR-SR measurements for B-doped and undoped SiNWs. A peak was observed at about 625 cm–1 for B-doped SiNWs, whereas no peak was observed for undoped SiNWs, as shown in Figure 4 (d). The intensity of the 625 cm–1 peak closely depends on the B concentration in B-doped SiNWs and increases with increasing B concentration [126], indicating a B vibrational peak. The peak position is very close to the local vibrational peak for 11B in SiNWs observed by Raman spectroscopy, but is slightly different. The difference in the peak position between IR-SR measurement and Raman spectroscopy measurement can be explained by the difference in measurement temperature: the lower the measurement temperature, the higher the wavenumber. In addition, the reason that the absorption peak for the 10B isotope peak is not observed in the IR-SR measurement is that the peak for 10B is about 1/4 that of the peak for 11B, putting it below the detection limit for this measurement due to the large background. When the sample after HAT treatment (180 ˚C for 30 min) was measured in the same way as with the Raman spectroscopic experiments, a peak was observed on the high wavenumber side, caused by hydrogen absorption.The electrical activation of B in Si observed by IR measurements can be evaluated by observing the peak of the electronic transition. Burstein et al. reported the first observation of the discrete electronic transition of bound holes from the ground state of neutral B acceptor atoms to a series of hydrogen-like excited levels below the band edge using low-temperature IR measurements of B-doped bulk Si [127].  Other groups later reported similar observations [128-131]. In their study of B-doped bulk Si, peaks due to electronic transitions of 1s-2p and 1s-3p were observed at 278 and 320 cm–1. To observe the electronic transition peaks for B-doped SiNWs, IR-SR measurements were performed in the far-infrared region, as shown in Figure 4 (f). Peaks can be seen at about 278 and 319 cm–1, which are almost the same as the peak positions observed with B-doped bulk Si. This result indicates that these peaks are due to the electronic transitions of 1s-2p and 1s-3p in B-doped SiNWs and that B is electrically activated in SiNWs. Observation of electronic transitions using IR measurements showed the transmittance of IR light by the sample to sharply decrease with increasing B concentration in bulk Si, preventing the electronic transition peak from being observed. On the other hand, since the NWs are very thin, IR light is easily transmitted, even at high B concentrations. The high brilliance of synchrotron radiation allows electron transitions to be detected even at a high B concentration of 1018 cm–3. The full width at half maximum (FWHM) of the electron transition peak increases with increasing B concentration, rising to about 7 cm–1 at a B concentration of 1018 cm–3.3.2 P doping in SiNWsHere we examine the results of P doping as an n-type dopant impurity [50-52,60]. We begin with an example of energy-dispersive X-ray spectroscopy (EDX) measurements performed during TEM observation in P-doped SiNWs. In Figure 5 (a), a peak due to Si is clearly observed at about 1.7 keV, and a small peak is observed at about 2 keV on the high energy side. This peak, caused by P, demonstrates that P is present in the SiNWs. The peak due to Mo is the signal from the TEM grid. Next, Raman spectroscopy, which was used for p-type B-doped SiNWs, is applied. Since the mass of P is much closer to that of Si than that of B, the localized vibrational peak of P cannot be observed because of Si’s large optical phonon peak. On the other hand, although the Fano effect used in the evaluation of electrical activity can be evaluated, because there are no clear levels such as heavy holes and light holes in the valence band in Si, the energetic interference with discrete levels of phonons weakens in n-type P-doped Si, and as a result, the asymmetry of the Si optical phonon peak due to the Fano effect shrinks. Figure 5 (b) shows the change in the Si optical phonon peak when P ion implantation is performed at room temperature and at 300 °C (1 x 1015 P+ cm–2). A result from before P ion implantation is also shown for comparison. The Si optical phonon peak is shifted to a lower wavenumber by P ion implantation and also shows slightly asymmetric broadening. This is the Fano effect in n-type Si. The lower wavenumber shift is more pronounced for ion implantation at 300 °C. This is due to the effect of hot implantation, which suppresses the internal polycrystallization of SiNWs by ion implantation and increases the electrical activity of P [56]. ESR is also used to evaluate the state and electrical activity of P in bulk Si. Like the IR method, ESR is normally used for bulk sample measurements and is therefore difficult to use directly to evaluate nanostructures. In the ESR method, the signal intensity increased even in the nanostructure as a result of adding SiNW powder to the microquartz capsule to enhance the detection sensitivity. The result is shown in Figure 5 (c). The signal observed at g = 2.005 in the absence of P doping is the ESR signal due to the Pb centers occurring at the oxide interface, as described in Section 2.2. A new ESR signal at a g value of 1.998 was observed when P-doping was applied during growth (in situ doping) and when P-ion implantation was performed [50-52,56,60]. This g value is known to be the signal from the conduction electrons of P in Si. The observation of ESR signals due to conduction electrons of P indicates that P atoms introduced into SiNWs locate the Si substitutional sites and are electrically activated at these sites. In other words, this is a result which proves that n-type SiNWs are formed. With hot implantation, the conduction electron signal intensity of P is increased because high-temperature ion implantation suppresses polycrystallization inside the SiNWs and increases the electrical activity of the doped P [56]. The electrical activity of P can be roughly estimated from the linewidth of the conduction electron signal of P. Figure 5 (d) shows the relationship between the linewidth and the carrier concentration. The linewidth of the conduction electron signal for P initially decreases with increasing carrier concentration. This is due to the motional narrowing and exchange narrowing effects of spins as the P concentration increases and the electron wave function overlaps. As the carrier concentration increases further, the interaction between spins becomes stronger, resulting in a shorter relaxation time and an increased linewidth. The results of Figure 5 (c) are plotted in Figure 5 (d) together with the values previously reported for bulk Si. This plot shows that high electrical activity in excess of 1019 cm–3 is obtained, even in SiNWs.3.3 B and P doping in GeNWsThe evaluation by Raman spectroscopy is also applicable to GeNWs. The peak observed at 300.2 cm–1 is the Ge optical phonon peak [64]. With B-doped GeNWs, the low-wavenumber shift of the Ge optical phonon peak and asymmetric broadening to the low-wavenumber side are clearly observed by B-doping (Figure 6 (a)). This change became marked when the excitation wavelength during Raman spectroscopy was lengthened. This result indicates that the observed downshift and asymmetric broadening of the phonon peak to the low-wavenumber side are due to the Fano effect. Two peaks were observed at 544 and 565 cm–1 (Figure 6 (b)). Each observed peak position corresponds to the position where the localized vibrational peaks for 11B and 10B in Ge are observed, and the interval between the peaks can be explained by this isotope shift [64]. The intensity ratios of the two peaks are consistent with the natural abundance of isotope B, except for the effect of another phonon-related peak at around 560 cm–1. Thus, the observed peaks are identified as the localized vibrational peaks for 11B and 10B in GeNWs. In the results for P doping, the local vibrational peak for P is about 345 cm–1 and asymmetric broadening to low wavenumbers due to the Fano effect is observed (Figure 6 (c) and (d)). The observations of the localized vibrational peak and the Fano effect demonstrate that B and P introduced into GeNWs enter the Ge substitution sites and are electrically activated, proving that p-type and n-type GeNWs are formed [52,64].4. Core-shell NWs4.1 Morphology, strain and defects in shell layersWe have so far described the growth, formation, doping and evaluation of single-structure SiNWs and GeNWs. However, as mentioned in the Introduction, carrier mobility is decreased due to scattering by the impurities doped into the NWs. As a structure which suppresses this effect, there is increasing research interest in core-shell type NWs in which a heterojunction is formed in the diameter direction inside the NW [78-98]. In the formation of core-shell NWs, the core NW can be formed by either top-down or bottom-up methods. After the formation of the core NW, the shell is formed by CVD [98]. It is important that the surface of the core NW is not oxidized: if the surface of the core NW is oxidized and shell formation is performed in an oxidized state, epitaxial growth is naturally inhibited.It is also important to optimize the shell formation conditions, even when there is no oxide film on the core NW surface. Since Si and Ge form heterojunctions, strain due to lattice mismatch occurs [84-86,90,91,95-98,134-138]. When the thickness of the shell layer exceeds the critical film thickness, the strain is dissipated by the appearance of misfit-dislocations and stacking faults due to lattice mismatching  (Figure 7 (a)). Depending on the formation conditions, Ge dots may be formed on the surface of the Si core NW rather than as a uniform shell layer [134-137]. This is the Stranski-Krastanow (S-K) growth mechanism, in which lower surface and edge energy, and energetic stabilization occur due to island growth in the presence of lattice mismatching [138-140]. Examples are shown in Figure 7 (b). 4.2 Hole gas accumulationIn the core-shell NWs composed of Si and Ge, the band structure is Type II, and hole gas can be formed in the Ge layer even when no impurity doping is performed. However, considering the high degree of integration of actual NW transistors and the operational control of a large number of integrated transistors, it is important to increase the carrier concentration. In this case, acceptor doping of the Si layer is needed for p-type channels. We have successfully observed the formation of i (intrinsic)-Ge/p-Si and p-Si/i-Ge core-shell NWs and the accumulation of hole gas in i-Ge core NWs and i-Ge shell layers. To realize the accumulation of hole gas in the i-Ge region, it is also important to suppress intermixing at the interface during the formation of the core-shell NWs. Although it is beneficial to form a highly crystalline shell layer at high temperatures, intermixing at the Si and Ge interfaces occurs under high-temperature conditions, and the steepness of the interface is impaired [141]. Further, when dopant impurity diffusion occurs during intermixing, the problem of impurity scattering recurs. When the core is Si, the temperature required for epitaxial growth of the Ge shell layer is 500 °C, at which intermixing does not occur. On the other hand, when the core is Ge, it is necessary to set the growth temperature above 700 °C for epitaxial growth of the Si shell layer and electrical activation of p-type impurities such as B. At this temperature, however, intermixing of Si and Ge may occur, although only near the interface, and interdiffusion of B atoms from the Si shell layer to the Ge core NWs may also occur as described below. The diffusion coefficient of Si in Ge at 700 °C is in the order of 10–15 to 10–16 cm2 s–1, while that of Ge in Si at 700 °C is much lower than 10–18 cm2 s–1 [142-144]. The diffusion of Ge atoms into Si is almost non-existent. Using the fastest diffusion coefficient, the maximum diffusion of Si atoms into the Ge layer was about 5 nm at 700 °C for 5 min. However, diffusion into the nanostructure is generally more difficult than in bulk (this is termed the self-purification effect), and Si and Ge intermixing appears to take place only at the interface of the Ge/Si core-shell NW. This intermixing may cause dopant diffusion, but only at the interface.Figure 8 (a) shows TEM and EDX images of a p-Si/i-Ge core-shell NW. The EDX image shows only the i-Ge shell layer. The core NW was formed by nanoimprint lithography and Bosch etching. Subsequently, the diameter of the Si NWs was reduced to about 50 nm by thermal oxidation and etching, and the i-Ge layer was grown by CVD after removing surface defects. The growth of i-Ge shell layers was performed using a GeH4 (100%) flux of 10 sccm at 500 °C. The total pressure was set at 8 Torr by mixing with N2 gas and the growth time was 60 sec. High-resolution TEM images of the i-Ge layer show that the shell layer has no defects and has a single crystal structure. The EDX image shows that the NW has a two-layer internal structure and that a core-shell structure is formed.The accumulation of hole gas in the i-Ge shell layer can be demonstrated by applying Raman spectroscopy. In core-shell nanowires doped with dopant impurities, stress due to heterojunctions is generated in the core and shell regions, so it is necessary to distinguish the effects of stress from those of doping. First, the results of Raman spectroscopy shown in Figure 8 (c) show Ge optical phonon peaks for bulk Ge and p-Si/i-Ge core-shell NWs grown with i-Ge layers for 80 seconds. In core-shell NWs, the Ge optical phonon peak shifts to the higher wavenumber side. This is because Ge, which has a larger lattice constant than Si, grows heteroepitaxially on the p-Si NW surface, and compressive stress is applied from the Si core NW to the i-Ge layer. With increasing shell formation time, the Ge optical phonon peak shifts to a lower wavenumber and shows a lower value than the optical phonon peak for bulk Ge (Figure 8 (d) and (e)). This is due to the Fano effect as described above. Since the Ge layer is not doped with impurities, this result indicates that hole gas is formed in the Ge layer due to the formation of the core-shell structure. Figure 8 (b) summarizes the changes in Raman shift in diagrammatic form.Hole gas accumulation in the i-Ge shell layer can be further demonstrated by constructing the following p-Si/i-Ge/p-Si core-able shell NWs. Figure 9 (a) shows an EDX image of the p-Si/i-Ge/p-Si core-able shell NWs. It can be seen that a three-layer heterojunction is formed inside the NW. Due to the formation of this structure, the accumulation of hole gas inside the i-Ge layer is enhanced by the effects of the p-Si shell. Figure 9 (c) shows the results of actual Raman measurements. Compared with the p-Si/i-Ge core-shell NWs, the Ge optical phonon peak shifts further to lower wavenumbers in p-Si/i-Ge/p-Si core-double shell NWs that have an additional p-Si shell layer. This low wavenumber shift also increases with increasing thickness of the p-Si shell layer. The Fano parameters, q and Γ, were obtained by fitting of the Fano equation. These results indicate that hole gas accumulation in the i-Ge shell layer is enhanced because the Fano effect is enhanced. Thus, by forming a core-shell structure, hole gas can be generated in a region where no impurities are present, and its concentration can be controlled by increasing the B concentration in Si, forming a double-shell structure. As described above, the core-shell NW structure makes it possible to separate the impurity-doped region from the carrier transport region, thereby realizing a high-mobility channel without impurity scattering.5. ConclusionWe conducted a comprehensive review of research on the formation of Group IV semiconductor NWs, impurity doping for functionalization, and development of high-mobility channels based on core-shell heterostructures. Research on the formation of NWs started using the bottom-up method employing the VLS growth mechanism, seeded with a metal catalyst. Subsequently, top-down techniques that enabled NW size and alignment control took over. However, the top-down approach causes damage to the NW surface and requires a process for removing the damaged layer.If NWs are used as channels in transistors, pn control by impurity doping is of key importance. Although it is necessary to evaluate the state and electrical activity of impurities introduced into the NWs, it is difficult to measure the trace amounts of impurities introduced into nanoscale NWs. Here, the results obtained by Raman spectroscopy that can be applied in a nondestructive and noncontact way, infrared absorption spectroscopy using synchrotron radiation, and electron spin resonance at low temperatures were mainly adopted. We couldn't go into more detail about the atomic probe method this time, however it is also utilized in the NW research field and makes it possible to evaluate the distribution and concentration of impurities. NWs doped with impurities face the problem that mobility decreases due to scattering by the introduced impurities. As a solution to this, we describe a structure designed to prevent impurity scattering and current research into core-shell NWs in which the impurity doped region and carrier transport region are separate. This structure is called an NW-HEMT, and it is the HEMT which is a focus of active research in compound semiconductors. In Raman spectroscopy, the Fano effect can be used to distinguish and evaluate the carrier region inside the heterojunction, and it can be proven that hole gas is generated in the undoped Ge region. Our results indicate that core-shell NWs can be used to form channels without impurity scattering.AcknowledgementsThis work was supported by JSPS Kakenhi (Nos. 21H04642 and 20K21135), The Murata Science Foundation and the World Premier International Research Center Initiative (WPI Initiative), MEXT, Japan.Data availability statement All data that supports the findings of this study is included within the article and any supplementary files.Conflict of interest The authors declare no competing interests. ORCID iDs Naoki Fukata https://orcid.org/0000-0002-0986-8485Wipakorn Jevasuwan https://orcid.org/0000-0001-9117-2497References[1] Fundamental Properties of Semiconductor Nanowires edited by Naoki Fukata and Riccardo Rurali (Springer Nature Singapore Pte Ltd. 2021).[2] Li Y, Qian F and Lieber C M 2006 Nanowire electronic and optoelectronic devices Mater. 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Phys. 71 1125-1171.[140] Politi P, Grenet G, Marty A, Ponchet A and Villain J 2000 Instabilities in crystal growth by atomic or molecular beams Phys. Rep. 324 271.[141] Zhang X, Jevasuwan W and Fukata N 2020 Interfacial intermixing of Ge/Si core-shell nanowires by thermal annealing Nanoscale 12 7572-7576.[142] Silvestri H H, Bracht H, Hansen J L, Larsen A N and Haller E E 2006 Diffusion of Silicon in Crystalline Germanium Semicond. Sci. Technol. 21 758–762.[143] Gavelle M, Bazizi E M, Scheid E, Fazzini P F, Cristiano F, Armand C, Lerch W, Paul S, Campidelli Y and Halimaoui A 2008 Detailed Investigation of Ge-Si Interdiffusion in the Full Range of Si1-xGex (0 ≤x≤1) Composition J. Appl. Phys.104 113524.[144] Castrillo P, Pinacho R, Jaraiz M and Rubio J E 2011 Physical Modeling and Implementation Scheme of Native Defect Diffusion and Interdiffusion in SiGe Heterostructures for Atomistic Process Simulation J. Appl. Phys. 109 103502.Figure 1. Schematic illustrations of SiNW formation by (a) catalytic laser ablation, (c) CVD and (e) nanoimprint lithography and the Bosch method. SEM images of SiNWs formed by (b) catalytic laser ablation, (d) CVD and (f) nanoimprint lithography and the Bosch method.Figure 2. SEM images of p-SiNW (a) before and after annealing in (b) Ar and (c) H2 gas atmosphere. (d) ESR signals observed for p-SiNW arrays before and after annealing in Ar and H2 gas atmosphere. (e) A model for surface morphology of SiNW annealed in Ar and H2 gas atmospheres. Reproduced from [106], © 2021, with the permission of IOP Publishing Ltd.Figure 3. SEM images of SiNW arrays (a) before and (b) after thermal oxidation and HF etching. Reproduced from [106], © 2021, with the permission of IOP Publishing Ltd.Figure 4. (a) Raman spectra of B-doped and undoped SiNWs. (b) Raman spectra of B-doped SiNWs before and after hydrogen atom treatment (HAT). (c) Dependence of Raman spectra of B-doped and undoped SiNWs on excitation wavelength during Raman spectroscopy. (d) IR absorbance spectra of B-doped and undoped SiNWs. (e) IR absorbance spectra of B-doped SiNWs before and after hydrogen atom treatment (HAT). (f) IR absorbance spectra of B-doped SiNWs observed in the far-infrared region. Reproduced from [59], © 2006, with the permission of the American Institute of Physics and reproduced from [126], © 2015, with the permission of the Royal Society of Chemistry.Figure 5. EDX spectra of P-doped SiNWs. (b) Raman spectra of intrinsic SiNWs (as grown) and SiNWs implanted with P+ ions at RT and 300 °C. Thermal annealing was performed at 900 °C after P+ ion implantation. (c) ESR signals observed for as-grown intrinsic SiNWs, in situ P-doped SiNWs, and P-doped SiNWs by P+ ion implantation at RT and at 300 °C. (d) Dependence of linewidth of ESR signal of conduction electrons on P donor concentration. Reproduced from [60], © 2007, with the permission of the American Institute of Physics and reproduced from [56], © 2012, with the permission of the American Chemical Society.Figure 6. Dependence of Raman spectra of (a) B-doped and (c) P-doped GeNWs on excitation wavelength during Raman spectroscopy. (b) B localized vibrational peak and (d) P localized vibrational peak observed for B-doped and P-doped GeNWs, respectively. Reproduced from [64], © 2012, with the permission of the American Chemical Society.Figure 7. TEM images of p-Si/i-Ge core-shell NWs grown under different growth conditions.Figure 8. (a) TEM and EDX images of p-Si/i-Ge core-shell NWs. (b) Dependence of Raman shift of Ge optical phonon peak on shell growth time. (c-e) Ge optical phonon peaks observed for bulk Ge and p-Si/i-Ge core-shell NWs with different shell growth times. Reproduced from [98], © 2019, with the permission of the American Chemical Society.Figure 9. (a) EDX image and (b) schematic illustration of p-Si/i-Ge/p-Si core-double shell NWs. (c) Raman shift of Ge optical phonon peak observed for bulk Ge, p-Si/i-Ge core-shell NWs and p-Si/i-Ge/p-Si core-double shell NWs. Dependence of (d) q and (e) Γon the outer p-Si shell layer. Reproduced from [98], © 2019, with the permission of the American Chemical Society.xxxx-xxxx/xx/xxxxxx1© xxxx IOP Publishing Ltd_1305372875.unknown