2019 Volume 60 Issue 7 Pages 1168-1176
Severe plastic deformation (SPD) has been widely studied in order to enhance the strength and ductility of metallic materials. Among various SPD processing techniques, high-pressure torsion (HPT) can be applied to various brittle materials including semiconductors. In this overview, we report on the HPT processing of Si, Ge, and compound semiconductor GaAs. When crystalline Si was subjected to HPT, metastable body-centered-cubic (bcc) Si-III and rhombohedral Si-XII as well as amorphous regions were formed. After annealing, Si-III and Si-XII reversely transformed to diamond-cubic Si-I. No appreciable photoluminescence (PL) peak was observed from the as-HPT processed samples while a broad PL peak originating from Si-I nanograins appeared after annealing. The electrical resistivity was increased just after compression without anvil rotation, but it decreased after HPT-processing because of the formation of semimetallic Si-III. In the case of Ge, metastable tetragonal Ge-III was formed by room-temperature HPT processing. A broad PL peak originating from diamond-cubic Ge-I nanograins was observed after annealing. The metastable bcc Ge-IV was observed in the cryogenic-HPT-processed samples. In the case of GaAs, no metastable phase was observed in the HPT-processed samples. A strong PL peak associated with the bandgap disappeared after HPT processing. An additional PL peak in the visible light region appeared after annealing. These results suggested that noble properties such as optical and electrical properties can be obtained by applying HPT processing to semiconductor materials.
Fig. 20 Schematic diagram of phase transformation and grain refinement mechanism of semiconductor materials by HPT processing.
Severe plastic deformation (SPD) has been widely studied for producing ultrafine grained metals with high strength and ductility.1–3) The application of SPD processing to semiconductor materials is of great interest to explore novel functionalities of the materials, because unique optical properties such as photoluminescence appear from nanocrystalline Si4,5) and Ge.6,7) The semiconductor nanocrystals are commonly formed by bottom-up techniques such as sputtering,5–7) chemical vapor deposition,8–10) and ion implantation,11,12) or top-down techniques such as electrochemical etching13,14) and ball milling.15) The SPD technique is one of the top-down techniques, and is possible to obtain a bulk form of semiconductor nanocrystals. In addition, allotropic phase transformation can occur during SPD processing due to high-pressure conditions. When high pressure is applied to crystalline Si, diamond-cubic Si-I transforms to Si-II having a β-Sn structure at ∼11 GPa, and then further transforms to various high-pressure phases such as orthorhombic Si-XI (∼13 GPa) and simple hexagonal Si-V (∼15 GPa).16,17) Upon unloading the pressure, it transforms not only back to Si-I but also to metastable phases such as rhombohedral Si-XII and body-centered-cubic Si-III.18) Since these metastable phases exhibit unique electrical and optical properties,19–22) the SPD-processed semiconductor materials may open up new device applications.
Among various SPD techniques, high-pressure torsion (HPT)23–25) is promising because it can be applied for various brittle materials such as intermetallics26) and ceramics.27–29) The HPT processing of semiconductor materials have been reported by several researchers. The first HPT processing to Si and Ge was reported by Bridgman.23) Blank and Kulnitskiy30) carried out the HPT processing of Si, and reported the direct transitions of Si-I to metastable phases under shear deformation. Islamgaliev et al. further reported the nanograin refinement and the formation of metastable phases of Ge31,32) and Si,33) and photoluminescence from the HPT-processed Si.33) However, the detailed crystal structural properties and functional properties such as electrical resistivities and photoluminescence have not fully understood so far. In this overview, we report the recent progress of HPT processing of group-IV semiconductors Si and Ge, and III-V compound semiconductor GaAs, and show their structural and functional properties of the HPT-processed and successive annealed samples.
HPT experiments were carried out using single crystalline Si(100) wafers, high-purity (99.999%) polycrystalline Ge, and single crystalline GaAs(100) wafers. They cut into ∼5 × 5 mm2 pieces or disks having a diameter of 5 mm, and put into the hole of the lower anvil with a depth of 0.25 mm and a diameter of 0.5 mm. The details of the HPT facility are described elsewhere.34) The HPT was operated at room temperature. For cryogenic HPT processing of Ge, the anvils and disk samples were immersed in liquid nitrogen. The temperature during cryogenic processing was monitored by a thermocouple located in the upper anvil. The loading pressure was set to 24 GPa. It should be noted that the loading pressure was the nominal value given by the applied load divided by the total area of the central shallow hole on the anvils. The rotation speed of the lower anvil was set to 1 rpm. The HPT-processed samples were basically polished to a mirror-like surface, and some of the samples were further annealed in nitrogen or argon atmosphere. The processed samples were characterized by x-ray diffraction (XRD) with the Kα radiation, micro-Raman and photoluminescence (PL) measurements with laser operating at a wavelength of 488 nm, and transmission electron microscopy (TEM). Electrical resistivities were measured by a four-point probe method with a probe distance of 1 mm. The core-level and valence band photoemission spectra were taken by synchrotron radiation using the beamline BL13 equipped with a planar-type undulator and a varied-line-spacing plane-grating monochromator, at the Saga Light Source.35)
The HPT-processed Si consists of Si-I and metastable phases as well as amorphous regions. Figure 1 shows XRD profiles of the HPT-processed Si. The XRD profile of the sample after mere compression (N = 0) shows that all the diffraction peaks correspond to Si-I. When the number of rotations (N) increases to 10, additional diffraction peaks especially at ∼35°, 53°, and 92° corresponding to metastable Si-III and Si-XII appear and peak broadening occurs in the Si-I peaks.36) The intensities of the Si-III/XII diffraction peaks increase with increasing the number of rotations as reported earlier.37) Figure 2 shows a micro-Raman spectrum of the HPT-processed sample. The Raman spectrum shows that peaks appear at ∼160, 380, 415, and 430 cm−1 for Si-III, and ∼350 and ∼396 cm−1 for Si-XII, other than a peak at 520 cm−1 for Si-I.38) Broad amorphous peaks also appear at ∼150 cm−1 and ∼470 cm−1.38) TEM observations revealed that the HPT-processed samples consist of micrograins and nanograins of Si-I and metastable phases.36,37,39) Figure 3 shows a typical high-resolution TEM (HRTEM) image and fast-Fourier transform (FFT) patterns of the HPT-processed Si. The inset FFT pattern (at top-right corner) shows polycrystalline rings and amorphous halo. The FFT patterns taken from the region a and the regions b and c in Fig. 3 correspond to Si-III and Si-I, respectively. A close observation reveals that metastable Si-IV having a hexagonal diamond or lonsdaleite structure is also observed in the HPT-processed sample.39) The HPT-processed samples were further annealed in nitrogen atmosphere. Figure 4 shows the XRD profile and Raman spectrum of the HPT-processed sample after annealing at 600°C for 1 hour. The diffraction peaks corresponding to Si-III and Si-XII disappear and only Si-I peaks are present in the XRD profile of Fig. 4(a). The Raman spectrum also shows that only a Si-I peak at ∼520 cm−1 is present as shown in Fig. 4(b).
XRD profiles of HPT-processed Si after compression (N = 0) and for N = 10.
Micro-Raman spectrum of HPT-processed Si for N = 10.
HRTEM image and FFT pattern (top right) of HPT-processed Si for N = 10. FFT patterns of a–c in insets are taken from regions indicated by square regions a–c.
(a) XRD profile and (b) micro-Raman spectrum of HPT-processed Si for N = 10 after annealing at 600°C for 1 hour.
Figure 5 shows PL spectra of the samples after HPT processing for N = 20 and after annealing at 600°C for 2 hours.37) No appreciable peak appears for the HPT-processed sample, while a broad peak is observed from the annealed sample.36,37) Since Si-I nanograins are present after the annealing at 600°C,39) the observed PL peak is due to quantum confinement in Si-I nanograins as observed in the case of nanocrystalline Si.5) No appreciable PL from the HPT-processed sample indicates the presence of a high density of lattice defects in the nanograins, which suppresses the quantum confinement.36) Figure 6 shows the resistivity changes with respect to the HPT processing with and without successive annealing of normally doped (n-Si), heavy doped (n+-Si), and ultraheavy doped (n++-Si) samples.40) The resistivity increases by 101–102 after mere compression. The resistivity decreases to ∼0.1 Ω cm for n- and n+-Si, and to ∼0.02 Ω cm for n++-Si after HPT processing for N=10 despite the grain size is refined through HPT processing.40) Such a resistivity decrease is opposite to the HPT-processed Cu,41) and can be explained by the appearance of metastable phases, especially semimetallic Si-III. Since the Rietveld analysis revealed that the volume fraction of Si-XII in the HPT-processed samples is less than 10%,40) the formation of Si-III in the HPT-processed samples results in a major contribution to the decrease in the electrical resistivity. The significant increase in the resistivity after the annealing at 600°C can be explained by the transformation of Si-III/XII to Si-I retaining nanograins.39)
PL spectra of HPT-processed (as-HPT) sample for N = 20 and after annealing at 600°C.37) Sharp luminescence peaks are due to laser plasma lines.
Resistivity changes with respect to HPT processing and successive annealing for N = 10 of n, n+, and n++-Si.40)
In order to investigate the electronic structure of the HPT-processed Si, we carried out photoemission spectroscopy measurements using synchrotron radiation.42) Figure 7 shows the Si core level (Si 2p) spectra of the HPT-processed p- and n-Si samples after Ne+ sputtering at 1 kV for 30 min and subsequent annealing. The spectra after sputtering show the Si–Si peak including the Si 2p3/2 at ∼99.2 eV and 2p1/2 spin orbit doublet components as well as a weak broad suboxide peak43) at 100.3–100.4 eV due to knocked-on oxygen from native oxide originating from sputtering. No shift appears for Si 2p3/2 peaks after annealing up to 200°C. Figure 8 shows valence band (VB) spectra obtained from the HPT-processed p- and n-Si samples after sputtering and subsequent annealing. The VB spectra after sputtering have peaks at ∼2.4 eV and ∼7 eV. These peaks correspond to the main peaks in the density of states for Si-III/XII (∼2.3 eV)44) and Si-I (∼7 eV).45) The gradual decrease of the peak at ∼2.4 eV with increasing the annealing temperature indicates the decrease of Si-III/Si-XII. Figure 9 shows detailed VB edge spectra of Fig. 8. The VB maxima are present at ∼0.3 eV after annealing, and shift by ∼0.4 eV after annealing. These results indicate that the bandgap of Si-III/XII in the HPT-processed samples is smaller than that of Si-I. It should be noted that Si-III transforms to Si-IV at 200°C,19,46) and that the bandgap of Si-IV is expected to be 0.95 eV.20) These results indicate that the HPT-processed samples consist of Si-I and metastable phases including Si-III and Si-XII, and suggest that the annealed samples consist of mixture of Si-I and Si-IV. It should also be noted that mid-gap pinning of the Fermi level also affects the VB maxima.
Si 2p photoemission spectra of HPT-processed p- and n-Si samples after sputtering and subsequent annealing. Photon energy was set at 130 eV. All spectra were taken at ∼1 mm from center of disk samples.
Valence band (VB) spectra obtained from HPT-processed p- and n-Si samples after sputtering and subsequent annealing. Photon energy was set at 130 eV. All spectra were taken at ∼1 mm from center of disk samples.
Detailed VB edge spectra of HPT-processed p- and n-Si samples after sputtering and subsequent annealing. Spectra were taken at same position of Fig. 8.
The HPT-processed Si also shows interesting thermal properties. The thermal conductivity measurements using picosecond time domain thermoreflectance show that the HPT-processed Si has a lattice thermal conductivity reduction by a factor of ∼20 due to the formation of nanograins and metastable Si-III/XII phases.47) This result indicates that the HPT-processed Si is promising for thermoelectric materials.
As for the case of Si, crystalline Ge exhibits a diamond cubic Ge-I at room temperature, and transforms to tetragonal Ge-II with the β-Sn structure at ∼10 GPa.48,49) Upon unloading, metastable Ge-III having a simple tetragonal structure is formed at room temperature.48–50) The formation of Ge-III is also observed in indentation experiments.38,51) When HPT processing is applied to crystalline Ge, not only grain refinement but also the formation of metastable phase are expected to occur. Figure 10 shows XRD profiles of the HPT-processed Ge samples.52) The diffraction peaks of the sample after compression (N = 0) correspond to Ge-I having a diamond-cubic structure. Additional peaks corresponding to simple tetragonal Ge-III appear from the HPT-processed samples for N = 5 and 10. Figure 11 shows a micro-Raman spectrum of the HPT-processed sample for N = 10. The Raman spectrum contains peaks at ∼151, 188, 215, 231, 248, and 276 cm−1. These peaks correspond to Ge-III.38,51) TEM observations revealed that the HPT-processed samples consists of micrograins and nanograins of Ge-I and Ge-III.52–54) Figure 12 shows an HRTEM image and the fast-Fourier transform (FFT) patterns obtained from the HPT-processed Ge.54) The size of the nanograins in Fig. 12 is ≤ 25 nm. The FFT pattern taken from the whole image of Fig. 12 consists of polycrystalline Ge-I and Ge-III rings as well as a weak amorphous halo. The FFT patterns taken from square regions indicate that Ge-I and Ge-III nanograins are formed in the HPT-processed sample. In addition to Ge-I and Ge-III nanograins as well as amorphous regions, highly defective regions including dislocations, nanotwins, and stacking faults are also observed in the HPT-processed samples.54)
Micro-Raman spectrum of HPT-processed Ge for N = 10.
HRTEM image and FFT pattern (top left) of HPT-processed Ge for N = 10.54) Insets are FFT patterns taken from regions indicated by squares.
The metastable Ge-III phase is reversely transformed to Ge-I after annealing or intense laser irradiation during Raman measurements. Figure 13 shows XRD profiles and micro-Raman spectra of the HPT-processed Ge after annealing at 300, 500 and 700°C.53) The Ge-III diffraction and Raman peaks disappear and only Ge-I peaks are visible after annealing. The intensity of the Ge-I peaks in XRD and Raman spectra become shaper and intense with increasing the annealing temperature. Figure 14 shows PL spectra of the HPT-processed Ge after annealing at 300, 500 and 700°C.53) A broad PL in the range of 600–800 nm appears after annealing at 300°C. Such a broad PL peak is also observed from the HPT-processed Ge after intense laser irradiation.53) The PL intensity decreases with increasing the annealing temperature. HRTEM observations revealed that Ge-I nanograins and residual Ge-III nanograins are present after annealing at 300°C.54) These results indicate that the observed PL originates from Ge-I nanograins due to quantum confinement effect.6,7)
(a) XRD profiles and (b) micro-Raman spectra of HPT-processed Ge for N = 10 after annealing at 300, 500 and 700°C.53)
PL spectra of HPT-processed Ge after annealing at 300, 500, and 700°C.53) Sharp luminescence peaks are due to laser plasma lines.
In order to investigate the temperature effect on the formation of metastable phases during HPT processing, we carried out HPT at cryogenic temperature. Figure 15 shows time-dependent XRD profiles of the samples processed at 100 K.52) The XRD profile of the sample just after HPT processing shows very weak body-centered-cubic Ge-IV peaks as well as the Ge-I peaks. The Ge-IV diffraction peaks still appear after keeping the sample at room temperature for 1 hour, but they disappear and only Ge-I peaks are present after 1 day. HRTEM observations revealed that Ge-I and residual Ge-III nanograins as well as amorphous phase are present in the cryogenic HPT-processed samples.54)
Time-dependent XRD profiles of cryogenic HPT-processed Ge for N = 10.52) XRD profiles were taken just after HPT and after keeping at room temperature for 1 hour and 1 day.
Based on the above results, the formation of Ge metastable phases processed by HPT processing can be explained as follows. The formation of Ge-III after HPT processing at room temperature for N = 5 and 10 as shown in Figs. 10 and 11 indicates that shear enhances the transformation from Ge-II to Ge-III.49,55) No appreciable Ge-IV phases are observed after room-temperature HPT processing because Ge-IV is observed by rapid pressure release (<1 s) from Ge-II.56) The Ge-IV is also obtained by high-pressure experiments at low temperature,57) but it is unstable at room temperature.58) The disappearance of Ge-IV in the cryogenic HPT-processed sample after keeping at room temperature as shown in Fig. 15 is due to annealing at room temperature which has been observed in the HPT-processed Cu at cryogenic temperature.59) These results indicate that shear strain as well as temperature during HPT processing takes an important role for the formation of Ge metastable phases.
Crystalline GaAs exhibits a direct bandgap of 1.42 eV at room temperature.60) The bandgap corresponds to near-infrared region while visible PL associated with quantum confinement effect is observed from nanocrystalline GaAs.61–63) When high pressure is applied to crystalline GaAs, zincblende GaAs-I transforms to orthorhombic GaAs-II at ∼17 GPa, and GaAs-II reversely transforms to GaAs-I via a cinnabar phase at 11–8 GPa upon unloading.64) In order to investigate the nanograined structure and optical properties of the bulk nanograined GaAs, we carried out HPT processing to crystalline GaAs at room temperature.65)
Figure 16 shows XRD profile of the HPT-processed GaAs for N = 0 and 10.65) All diffraction peaks correspond to zincblende GaAs-I. The sharp and intense diffraction peaks are observed in the XRD profile for N = 0, while peak broadening occurs due to grain refinement when N is increased to 10. Figure 17 shows micro-Raman spectra of the HPT-processed GaAs.65) The Raman spectrum of GaAs(100) shows a strong longitudinal optical (LO) phonon peak at 291 cm−1.66) A transverse optical (TO) phonon peak at ∼267 cm−1 in addition to the LO peak appears in the Raman spectra of the HPT-processed GaAs. The intensity of the LO peak decreases and that of the TO peak increases with increasing the distance from the sample center. The LO and TO peaks become broader and gradually shift to lower wave numbers as the distance from the sample center increases. Such peak broadening and peak shifts toward lower wavenumbers are observed in the case of GaAs nanocrystals.67) These results indicate that the grain refinement is enhanced by imposed shear strain. In fact, TEM observations reveal that micrograins (∼5 µm) and nanograins (∼10 nm) are observed in the HPT-processed GaAs.65)
XRD profiles of HPT-processed GaAs for N = 0 and 10.65) LaB6 peaks were used for XRD calibration.
Micro-Raman spectra of HPT-processed GaAs for N = 10.65) r denotes distance from disk center.
The HPT-processed GaAs samples were further annealed in Ar atmosphere. Figure 18 shows micro-Raman spectra of the HPT-processed GaAs samples after annealing. In addition to the original TO and LO phonon peaks of GaAs-I, As-related peaks appear at ∼200 and 260 cm−1 for the Eg and A1g phonons.68) The appearance of As-related peaks after the annealing indicates that the grain boundaries introduced by HPT processing play an important role in As precipitations. The XRD measurements revealed that all diffraction peaks of the annealed samples correspond to GaAs-I, and become sharp and intense due to grain coarsening.65) Figure 19 shows PL spectra of the HPT-processed GaAs and after annealing. No PL peak is observed in the HPT-processed sample because of a high density of lattice defects such as dislocations.69) A weak and broad peak around 700 nm appears after annealing at 400°C. A strong PL peak corresponding to bulk GaAs becomes prominent after annealing at 600°C. It should be noted that the broad PL peak around 700 nm is not related to oxides such as Ga2O3, As2O3, and As2O5, because they have broad PL peaks in the UV and blue ranges for Ga2O3,70) and green region for As2O3, and As2O5.71) The micro-Raman spectra of Fig. 18 also indicate that no appreciable oxide peaks appear in the annealed samples. It should also be noted that micrograins and nanograins still remain in the HPT-processed samples after annealing at 500°C.65) These results indicate that the broad PL peak around 700 nm originates from GaAs-I nanograins with less lattice defects due to quantum confinement effect.
Micro-Raman spectra of HPT-processed samples after annealing at 400°C and 600°C. All spectra were taken at 1 mm from center of disk samples.
PL spectra of HPT-processed GaAs for N = 10 and successive annealing at 400 and 600°C. All spectra were taken at 1 mm from center of disk samples. Sharp luminescence peaks are due to laser plasma lines.
No metastable phases have been observed in neither HPT-processed GaAs nor successive annealed samples as shown in Figs. 16–18. McMahon et al. reported that an SC16 phase with simple cubic structure is observed by heating the GaAs-II to above ∼127°C around 14 GPa, and is stable at atmospheric pressure.72) It is indicated that the frictional heating during HPT processing is negligible, and is estimated to be below 127°C.
The phase transformation and grain refinement mechanism of semiconductor materials by HPT processing are summarized in Fig. 20. When high pressure is applied to crystalline semiconductor materials, they transform to metallic high-pressure phases. When the rotation is applied to the sample under high pressure, grain refinement occurs due to a high density of lattice defects by torsional strain. Metastable phases such as Si-III/XII, and Ge-III/IV are formed upon unloading. These metastable phases transformed back to original Si-I and Ge-I by annealing. It is found that the formation of metastable phases, especially semimetallic Si-III affects the electrical and electronic properties, such as electric resistivity decrease and shift of the valence band. No appreciable PL was observed from the HPT-processed samples while a broad PL peak originating from Si-I, Ge-I, and GaAs-I nanograins with less lattice defects appeared after annealing. These results suggest that the HPT processing to semiconductor materials is promising to develop noble electrical and optical properties.
Schematic diagram of phase transformation and grain refinement mechanism of semiconductor materials by HPT processing.
The author is very much grateful to Prof. Zenji Horita of Kyushu University for giving me the opportunity to study HPT processing of semiconductor materials. The author would like to thank Dr. Kaveh Edalati of Kyushu University for HPT processing and valuable discussion, and Prof. Masamichi Kohno of Kyushu University and Prof. Junichiro Shiomi of University of Tokyo for resistivity measurements and for providing n++-Si wafers. The author would also like to thank Dr. Katsuhiko Saito and Prof. Qixin Guo of Saga University for Raman and PL measurements, Prof. Kazutoshi Takahashi of Saga University for photoemission experiments, and Prof. Martha R. McCartney and Prof. David J. Smith of Arizona State University for HRTEM observations. The author acknowledges the facilities for HPT in the International Research Center on Giant Straining for Advanced Materials (IRC-GSAM) at Kyushu University. The author also acknowledges the use of facilities in the John M. Cowley Center for High Resolution Electron Microscopy at Arizona State University. The photoemission experiments were performed at Saga University Beamline (SAGA-LS/BL13) with a proposal of H28-110V under the support of the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan. This work was supported in part by Grant-in-Aid for Scientific Research (S) (Grant No. JP26220909) from the Japan Society for the Promotion of Science, and in part by a Grant-in-Aid for Scientific Research in Innovative Areas “Bulk Nanostructured Metals” (Nos. JP22102004, JP25102708) from MEXT, Japan.