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Structural and electrical properties of high-quality 0.41 μm-thick InSb films grown on GaAs (1 0 0) substrate with InxAl1−xSb continuously graded buffer
Materials Research Bulletin 47 (2012) 2927–2930
Contents lists available at SciVerse ScienceDirect
Materials Research Bulletin journal homepage: www.elsevier.com/locate/matresbu
Structural and electrical properties of high-quality 0.41 mm-thick InSb films grown on GaAs (1 0 0) substrate with InxAl1 xSb continuously graded buffer Sang Hoon Shin a,b, Jin Dong Song a,*, Ju Young Lim a, Hyun Cheol Koo a, Tae Geun Kim b a b
Division of Future Convergence Technology, Korea Institute of Science and Technology, Seoul 136-791, Republic of Korea School of Electrical Engineering, Korea University, Seoul 136-701, Republic of Korea
A R T I C L E I N F O
A B S T R A C T
Article history: Available online 26 April 2012
High-quality InSb was grown on a GaAs (1 0 0) substrate with an InAlSb continuously graded buffer (CGB). The temperatures of In, Al K-cells and substrate were modified during the growth of InAlSb CGB. The cross-section TEM image reveals that the defects due to lattice-mismatch disappear near lateral structures in CGB. The measured electron mobility of 0.41 mm-thick InSb was 46,300 cm2/Vs at 300 K. These data surpass the electron mobility of state-of-the-art InSb grown by other methods with similar thickness of InSb. Crown Copyright ß 2012 Published by Elsevier Ltd. All rights reserved.
Keywords: A. Semiconductors A. Thin films B. Epitaxial growth D. Defects D. Electrical properties
1. Introduction InSb is a unique III–V compound semiconductor suitable for application in high-performance micro-Hall sensors [1], highspeed magnetic sensors [2,3], long-wavelength photodetectors [4– 6], and high-speed electronic devices [7], because of its ultra-high 300 K-electron mobility (m300 K) of >80,000 cm2/Vs [8] and narrow band-gap of 0.17 eV. Although InSb substrates are supplied by commercial vendors using the Czochralski method [9], the lack of semi-insulating substrates that are lattice-matched to InSb is a critical issue for practical applications, since the lattice-constant of InSb (6.48 A˚) is larger than readily available semi-insulating substrates such as GaAs (14% of lattice mismatch), InP (11%) and Si (18%). Consequently, a specific growth procedure has been developed to accommodate this large lattice mismatch. If widely studied GaAs substrates are chosen, InSb films are grown on them with sophisticated control of growth parameters [10,11], or with additional buffers including the use of compliant universal substrate [12], low temperature–high temperature (LT–HT) method [13,14], and a step-graded buffer (SGB) consisting of AlSb and InAlSb [10]. Finally, the growth of high-quality InSb on GaAs with a value of m300 K over 50,000 cm2/Vs has been achieved for a thickness of InSb larger than 2 mm [11,13,15]. However, even in this high-quality InSb films grown on GaAs and other substrates, defects such as threading dislocations (TDs) and micro twins (MTs) have been identified from transmission electron microscopy (TEM) images in the density of 13,000/cm [16]. Although this range of
* Corresponding author. Fax: +82 2 958 6851. E-mail address: [email protected] (J.D. Song).
defect density may be affordable in Hall sensors, it is clear that reducing the defect density is beneficial for the application to highspeed electron devices [17,18]. Moreover, for the isolation of devices and increased performance, the thickness of the InSbactive film should be reduced with increased quality. Sato et al. introduced 1 mm-thick InxAl1 xSb SGB by 9 steps of grading of compositions (from x = 0.1 to 0.9) on a 1 mm-thick AlSb buffer/ GaAs substrate and obtained a high-quality 2 mm-thick InSb active structure on it (m300 K 55,000 cm2/Vs) [15]. Meanwhile, we are expecting an inconveniently laborious process and a longer process time as a result of temperature changes of group-III Kcells for the sequential growths of step-graded composition of InAlSb. In this work, we have studied the characteristics of InSb film on GaAs (1 0 0) subtract by MBE growth and InAlSb continuously graded buffers (CGB) to alleviate lattice-mismatch between GaAs and InSb films. The m300 K of a 0.41 mm-thick InSb film on CGB turned out to be 46,300 cm2/Vs, which is comparable to that of SBG. 83% of the quality was achieved with 20% of the InSb thickness through a simpler method. 2. Experiment The samples used in this work were prepared on semiinsulating GaAs (0 0 1) substrates in a Riber compact 21E solid source molecular beam epitaxy (MBE) system. The surface oxide was removed from the GaAs substrate by heating the sample at substrate temperature (Ts) = 620 8C under the As2 mode. A 200 nmthick GaAs buffer layer was consecutively grown at Ts = 580 8C. The decrease of Ts to 450 8C was followed by the exchange of group V from As2 to Sb2. During the exchange, the high energy electron
0025-5408/$ – see front matter . Crown Copyright ß 2012 Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.materresbull.2012.04.121
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S.H. Shin et al. / Materials Research Bulletin 47 (2012) 2927–2930
The XRD measurement was performed by using a Bruker D-8 advanced diffractometer with Cu Ka radiation. The SIMS measurement was carried out by using IMS-4FE7 (CAMECA) with Cs+ ions. The morphological characteristics of the surface and the crosssectional sample structure were analyzed by AFM (PSIA) and TEM operating at 300 keV (Monochromatic and Cs corrected probe Titan 80-300), respectively. The van der Pauw Hall effect measurements were carried out with a 80 mm 64 mm pattern in the temperature range of 10–300 K by using a physical property measurement system. 3. Results and discussion
Fig. 1. Schematics of growth method of the sample.
diffraction (RHEED) pattern of (2 4) from the GaAs surface changed into (1 3) under the exposal of Sb2 [19]. Fig. 1 shows a schematic illustration of the sample structure and growth sequences. In the first phase, the flux of Al, which is equivalent to the growth rate of AlAs, and the Ts were fixed as 1.45 A˚/s (Al cell temperature; 1130 8C) and 450 8C, respectively. Meanwhile, the flux of In, which is equivalent to the growth rate of InP, increased from 0 (In cell temperature; 600 8C) to 3.80 A˚/s (755 8C) during 1 h. In the second phase, the flux of In was kept as 3.80 A˚/s. The Al cell temperature decreased to idle temperature (850 8C), and the Ts decreases to 400 8C within 1 h. During these two phases, the shutters of In, Al, and Sb2 were kept open. Finally, a 20 nm-thick In0.3Al0.7Sb cap layer was grown on it. It should be pointed out that the flux of Al from the Al cell was undetectable – considered as 0 A˚/s – around the Al cell temperature of 950 8C. Therefore, the corresponding flux of Al will be 0 A˚/s in the middle of the second phase. The beam equivalent pressure of the As2 flux was 3 10 6 Torr and the beam equivalent pressure of the Sb2 flux was 1.7 10 6 Torr.
Fig. 2(a) is the {1 1 1}, two-beam bright-field TEM image across the sample. The thickness of the film including the InAlSb capping layer, InSb and InAlSb CGB is found to be 2.39 mm. As shown in the image, in the vicinity of the InAlSb CGB/GaAs interface, large numbers of misfit dislocations along with neighboring strain fields resulting from lattice mismatch are observed. It is also found that as the distance increases from the CGB/GaAs interface, the density of misfit dislocations reduces gradually and finally the crystal defects including misfit dislocations, TDs, MTs, and stacking faults caused by lattice mismatch disappear around the distance of 2.0 mm from the CGB/GaAs interface. The reduction and disappearance of misfit dislocations with increasing distance from the GaAs substrate are also observed in the direct growth of InSb on GaAs. However, MTs and TDs remain in it even at a larger distance (>2 mm) where misfit dislocations disappear [10]. Meanwhile, in Fig. 2(a), neither a MT nor a TD is found around distances of over 2 mm. Because the density of the defects depends on the measurement positions where the crosssection TEM image was taken, it cannot be argued that the upper of InSb on CGB is free of defects. It is, however, noteworthy that the defects were effectively removed, and there is a mechanism to explain it. A similar reduction of crystal defects was also observed in InSb on InAlSb SGB and is attributed to a reduction of crystal defects at each interface of the SGB [15]. Mishima et al. [18] reported a TD filtering mechanism caused by the reflection of a TD between the abrupt interfaces of a higher Al-content Al0.24In0.76Sb layer and a lower Al-content Al0.12In0.88Sb layer.
Fig. 2. (a) {1 1 1}, two-beam bright-field TEM image across the sample. Eye-guide lines are inserted in the insets. (b) Estimated compositional profiles of InAlSb, as a function of depth. In-set shows normalized intensity measured by SIMS. (c) Comparison of measured and simulated XRD rocking curves.
S.H. Shin et al. / Materials Research Bulletin 47 (2012) 2927–2930
In the present study, the CGB structure does not have abrupt interfaces to filter MTs and TD. However, it is observed that lateral structures (LS) parallel to the GaAs interface are identified in the CGB region of Fig. 2(a). They are distorted due to strain fields, located below 2 mm of distance from the CGB/GaAs interface and scattered around the CGB. The origin of this structure is not clear but can possibly be attributed to the irregular relaxation of stress during CGB growth. However, as shown in the white circle of Fig. 2(a), a LS filters a TD. Consequently, the authors can apply the idea of TD filtering by LS [18] to understand the mechanism of reduction of the TD in the CGB. The distortion of TDs when bent and the resulting reflection of a TD is found in the blue circle of Fig. 2(a). This may also explain the reduction of defects. For a more in depth study of the compositional distribution of InAlSb CGB and thickness of the InSb, a SIMS measurement was taken. As shown in the inset of Fig. 2(b), In and Al atom counts were measured as a function of Cs+ ion sputtering time. The sputtering time when In or Al atom counts approach 0/s was considered as the CGB/GaAs interface. The composition x of InxAl1 xSb was calculated from the normalized intensity of In and Al atom counts and is illustrated in Fig. 2(b). The composition of InAlSb appears to be concave rather than linear, although the In and Al cell temperatures were linearly controlled by an MBE process computer software. This is reasonable considering the fact that flux of the source is not propositional to the source cell temperature. The abrupt decrease of x near the sample surface is due to the 20 nm-thick In0.3Al0.7Sb cap layer. The thickness of InSb was estimated as 0.41 mm where normalized Al atom counts are less than 0.01, and the composition x of InxAl1 xSb is estimated as less than 0.01 (see red arrow in Fig. 2(b) and its inset). From this estimation, it is found that the time when the flux of the Al approach 0 A˚/s is 42 min in the second phase, and the Al cell temperature for no flux out of the cell is calculated as 940 8C, which is in agreement with the measured Al cell temperature for no flux (950 8C) considering errors due to background noise in the flux sensor. Although this can be seen as an indirect support for SIMS measurements, an XRD measurement of the sample and an XRD simulation of the compositional information from the SIMS were compared for a more direct proof of the accuracy of the compositional information measured by the SIMS. As shown in Fig. 2(c), the measured rocking curve is in good agreement with the simulated rocking curves. With the SIMS (Fig. 2(b)) and XRD (Fig. 2(c)) information, the thickness of InSb can be identified as illustrated in Fig. 2(a). At this point, it is noteworthy that crystal defects disappear within InAlSb CGB and the structural quality of InSb is very high without the defects in Fig. 2(a). Meanwhile, dislocations are found in the crosssectional TEM image of the lower layer of InSb on SGB [15]. Fig. 3(a) shows the morphological information of 0.41 mm-thick InSb on InAlSb CGB with a 20 nm-thick InAlSb capping layer measured by AFM. The root-mean-square surface (RMS) roughness from a 10 mm 10 mm AFM image is 5.3 nm and the peak wavelength of a 2-dimensional (2-D) power spectral density (PSD) analysis is 5 mm. Compared with those of the SGB [10], the RMS roughness is similar to that of the 2 mm-thick InSb on a 500 nmthick InAlSb SGB/1 mm-thick AlSb/GaAs, and the peak wavelength is similar to that of the 2 mm-thick InSb on no InAlSb SGB/1 mmthick AlSb/GaAs. Considering that a larger mosaicity was found in thicker SGB with the larger RMS roughness and the larger peak wavelength of 2D-PSD [15] and that the peak wavelength of PSD is enlarged with increasing thickness of InSb [20], a smaller mosaicity may be expected in CGB due to the decreased thickness of InSb. Fig. 3(b) illustrates electron mobility as a function of measurement temperature. The dependence of mobility drops in a sigmoidal shape. These trends of temperature dependences are
2929
Fig. 3. (a) Surface 2-dimensional power spectral density analysis and AFM image (10 mm 10 mm) of the sample (inset). (b) Measured electron mobility as a function of measurement temperature and comparison of electron mobility at 300 K with references (inset).
in general accordance with the results for undoped InSb layers grown on GaAs [10,20]. The apparent rapid-drop of electron mobility with decreasing temperature in the low temperature range (80–270 K) can be interpreted as an artifact due to the presence of holes unintentionally caused by a p-type background [20]. The measured m300 K is 46,300 cm2/Vs. Finally, the m300 K of 0.41 mm-thick InSb on InAlSb CGB/GaAs is compared with various thicknesses of InSb on GaAs in the inset of Fig. 3(b) [11,13,15]. As shown there, the m300 K of InSb is explicitly enhanced, even under the assumption that the thickness of InSb is enlarged to 0.6 mm by 50% of error during the estimation of InSb thickness. This is attributed to defect reduction by InAlSb CGB. It should be pointed out that for increased high-speed performance, the thickness of InSb active film should be reduced with increased quality (m300 K). 4. Conclusions An InAlSb CGB was introduced to accommodate a lattice mismatch between GaAs (1 0 0) substrate and InSb. The defects caused by the lattice mismatch were effectively reduced at 2 mm from the interfaces between GaAs substrate and InAlSb CGB. This reduction of defect is attributed to LSs parallel to the GaAs interface which reflect, bend, stop and, as result, filter the TDs. The origin of this structure is not clear, however, can possibly be attributed to the irregular relaxation of stress during CGB growth. The thickness of pure InSb layers was found to be 0.41 mm, which was evaluated from SIMS and XRD measurements. The m300 K of 0.41 mm-thick InSb on top of InAlSb CGB is 46,300 cm2/Vs, which surpasses the m300 K of state-of-the-art InSb grown by other methods with
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similar thickness. We expect that CGB will be useful for high-speed device applications using InSb 2DEGs. Acknowledgments This work was supported mainly by KIST Institutional Programs including Dream Project and partially by NRF grant funded by MEST (No. 2011-0016471). References [1] A. Oral, et al. IEEE Trans. Magn. 38 (2002) 2438. [2] S.A. Solin, et al. Appl. Phys. Lett. 80 (2002) 4012. [3] S.A. Solin, et al. Science 289 (2000) 1530.
[4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20]
E. Michel, et al. IEEE Photonics Technol. Lett. 8 (1996) 673. W. Dobbelaere, et al. Appl. Phys. Lett. 60 (1992) 29. S. Ozer, C. Besikci, J. Phys. D: Appl. Phys. 36 (2003) 559. B. Bennett, et al. Solid State Electron. 49 (2005) 1875. A. Joullie´, et al. Mater. Res. Bull. 9 (1974) 241. Z. Burshtein, M. Azulay, Mater. Res. Bull. 19 (1984) 49. J. So¨derstro¨m, et al. Semicond. Sci. Technol. 7 (1992) 337. E. Michel, et al. Appl. Phys. Lett. 65 (1994) 3338. F.E. Ejeckam, et al. Appl. Phys. Lett. 71 (1997) 776. T. Zhang, et al. Appl. Phys. Lett. 84 (2004) 4463. M.C. Debnath, et al. J. Cryst. Growth 267 (2004) 17. T. Sato, et al. Physica E 21 (2004) 615. A. Okamoto, et al. J. Cryst. Growth 227 (2001) 619. T.D. Mishima, et al. Physica B 376 (2006) 591. T.D. Mishima, et al. Appl. Phys. Lett. 88 (2006) 191908. J.Y. Lim, et al. J. Korean Phys. Soc. 55 (2009) 1525. A. Okamoto, I. Shibasaki, J. Cryst. Growth 251 (2003) 560.
Contents lists available at SciVerse ScienceDirect
Materials Research Bulletin journal homepage: www.elsevier.com/locate/matresbu
Structural and electrical properties of high-quality 0.41 mm-thick InSb films grown on GaAs (1 0 0) substrate with InxAl1 xSb continuously graded buffer Sang Hoon Shin a,b, Jin Dong Song a,*, Ju Young Lim a, Hyun Cheol Koo a, Tae Geun Kim b a b
Division of Future Convergence Technology, Korea Institute of Science and Technology, Seoul 136-791, Republic of Korea School of Electrical Engineering, Korea University, Seoul 136-701, Republic of Korea
A R T I C L E I N F O
A B S T R A C T
Article history: Available online 26 April 2012
High-quality InSb was grown on a GaAs (1 0 0) substrate with an InAlSb continuously graded buffer (CGB). The temperatures of In, Al K-cells and substrate were modified during the growth of InAlSb CGB. The cross-section TEM image reveals that the defects due to lattice-mismatch disappear near lateral structures in CGB. The measured electron mobility of 0.41 mm-thick InSb was 46,300 cm2/Vs at 300 K. These data surpass the electron mobility of state-of-the-art InSb grown by other methods with similar thickness of InSb. Crown Copyright ß 2012 Published by Elsevier Ltd. All rights reserved.
Keywords: A. Semiconductors A. Thin films B. Epitaxial growth D. Defects D. Electrical properties
1. Introduction InSb is a unique III–V compound semiconductor suitable for application in high-performance micro-Hall sensors [1], highspeed magnetic sensors [2,3], long-wavelength photodetectors [4– 6], and high-speed electronic devices [7], because of its ultra-high 300 K-electron mobility (m300 K) of >80,000 cm2/Vs [8] and narrow band-gap of 0.17 eV. Although InSb substrates are supplied by commercial vendors using the Czochralski method [9], the lack of semi-insulating substrates that are lattice-matched to InSb is a critical issue for practical applications, since the lattice-constant of InSb (6.48 A˚) is larger than readily available semi-insulating substrates such as GaAs (14% of lattice mismatch), InP (11%) and Si (18%). Consequently, a specific growth procedure has been developed to accommodate this large lattice mismatch. If widely studied GaAs substrates are chosen, InSb films are grown on them with sophisticated control of growth parameters [10,11], or with additional buffers including the use of compliant universal substrate [12], low temperature–high temperature (LT–HT) method [13,14], and a step-graded buffer (SGB) consisting of AlSb and InAlSb [10]. Finally, the growth of high-quality InSb on GaAs with a value of m300 K over 50,000 cm2/Vs has been achieved for a thickness of InSb larger than 2 mm [11,13,15]. However, even in this high-quality InSb films grown on GaAs and other substrates, defects such as threading dislocations (TDs) and micro twins (MTs) have been identified from transmission electron microscopy (TEM) images in the density of 13,000/cm [16]. Although this range of
* Corresponding author. Fax: +82 2 958 6851. E-mail address: [email protected] (J.D. Song).
defect density may be affordable in Hall sensors, it is clear that reducing the defect density is beneficial for the application to highspeed electron devices [17,18]. Moreover, for the isolation of devices and increased performance, the thickness of the InSbactive film should be reduced with increased quality. Sato et al. introduced 1 mm-thick InxAl1 xSb SGB by 9 steps of grading of compositions (from x = 0.1 to 0.9) on a 1 mm-thick AlSb buffer/ GaAs substrate and obtained a high-quality 2 mm-thick InSb active structure on it (m300 K 55,000 cm2/Vs) [15]. Meanwhile, we are expecting an inconveniently laborious process and a longer process time as a result of temperature changes of group-III Kcells for the sequential growths of step-graded composition of InAlSb. In this work, we have studied the characteristics of InSb film on GaAs (1 0 0) subtract by MBE growth and InAlSb continuously graded buffers (CGB) to alleviate lattice-mismatch between GaAs and InSb films. The m300 K of a 0.41 mm-thick InSb film on CGB turned out to be 46,300 cm2/Vs, which is comparable to that of SBG. 83% of the quality was achieved with 20% of the InSb thickness through a simpler method. 2. Experiment The samples used in this work were prepared on semiinsulating GaAs (0 0 1) substrates in a Riber compact 21E solid source molecular beam epitaxy (MBE) system. The surface oxide was removed from the GaAs substrate by heating the sample at substrate temperature (Ts) = 620 8C under the As2 mode. A 200 nmthick GaAs buffer layer was consecutively grown at Ts = 580 8C. The decrease of Ts to 450 8C was followed by the exchange of group V from As2 to Sb2. During the exchange, the high energy electron
0025-5408/$ – see front matter . Crown Copyright ß 2012 Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.materresbull.2012.04.121
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S.H. Shin et al. / Materials Research Bulletin 47 (2012) 2927–2930
The XRD measurement was performed by using a Bruker D-8 advanced diffractometer with Cu Ka radiation. The SIMS measurement was carried out by using IMS-4FE7 (CAMECA) with Cs+ ions. The morphological characteristics of the surface and the crosssectional sample structure were analyzed by AFM (PSIA) and TEM operating at 300 keV (Monochromatic and Cs corrected probe Titan 80-300), respectively. The van der Pauw Hall effect measurements were carried out with a 80 mm 64 mm pattern in the temperature range of 10–300 K by using a physical property measurement system. 3. Results and discussion
Fig. 1. Schematics of growth method of the sample.
diffraction (RHEED) pattern of (2 4) from the GaAs surface changed into (1 3) under the exposal of Sb2 [19]. Fig. 1 shows a schematic illustration of the sample structure and growth sequences. In the first phase, the flux of Al, which is equivalent to the growth rate of AlAs, and the Ts were fixed as 1.45 A˚/s (Al cell temperature; 1130 8C) and 450 8C, respectively. Meanwhile, the flux of In, which is equivalent to the growth rate of InP, increased from 0 (In cell temperature; 600 8C) to 3.80 A˚/s (755 8C) during 1 h. In the second phase, the flux of In was kept as 3.80 A˚/s. The Al cell temperature decreased to idle temperature (850 8C), and the Ts decreases to 400 8C within 1 h. During these two phases, the shutters of In, Al, and Sb2 were kept open. Finally, a 20 nm-thick In0.3Al0.7Sb cap layer was grown on it. It should be pointed out that the flux of Al from the Al cell was undetectable – considered as 0 A˚/s – around the Al cell temperature of 950 8C. Therefore, the corresponding flux of Al will be 0 A˚/s in the middle of the second phase. The beam equivalent pressure of the As2 flux was 3 10 6 Torr and the beam equivalent pressure of the Sb2 flux was 1.7 10 6 Torr.
Fig. 2(a) is the {1 1 1}, two-beam bright-field TEM image across the sample. The thickness of the film including the InAlSb capping layer, InSb and InAlSb CGB is found to be 2.39 mm. As shown in the image, in the vicinity of the InAlSb CGB/GaAs interface, large numbers of misfit dislocations along with neighboring strain fields resulting from lattice mismatch are observed. It is also found that as the distance increases from the CGB/GaAs interface, the density of misfit dislocations reduces gradually and finally the crystal defects including misfit dislocations, TDs, MTs, and stacking faults caused by lattice mismatch disappear around the distance of 2.0 mm from the CGB/GaAs interface. The reduction and disappearance of misfit dislocations with increasing distance from the GaAs substrate are also observed in the direct growth of InSb on GaAs. However, MTs and TDs remain in it even at a larger distance (>2 mm) where misfit dislocations disappear [10]. Meanwhile, in Fig. 2(a), neither a MT nor a TD is found around distances of over 2 mm. Because the density of the defects depends on the measurement positions where the crosssection TEM image was taken, it cannot be argued that the upper of InSb on CGB is free of defects. It is, however, noteworthy that the defects were effectively removed, and there is a mechanism to explain it. A similar reduction of crystal defects was also observed in InSb on InAlSb SGB and is attributed to a reduction of crystal defects at each interface of the SGB [15]. Mishima et al. [18] reported a TD filtering mechanism caused by the reflection of a TD between the abrupt interfaces of a higher Al-content Al0.24In0.76Sb layer and a lower Al-content Al0.12In0.88Sb layer.
Fig. 2. (a) {1 1 1}, two-beam bright-field TEM image across the sample. Eye-guide lines are inserted in the insets. (b) Estimated compositional profiles of InAlSb, as a function of depth. In-set shows normalized intensity measured by SIMS. (c) Comparison of measured and simulated XRD rocking curves.
S.H. Shin et al. / Materials Research Bulletin 47 (2012) 2927–2930
In the present study, the CGB structure does not have abrupt interfaces to filter MTs and TD. However, it is observed that lateral structures (LS) parallel to the GaAs interface are identified in the CGB region of Fig. 2(a). They are distorted due to strain fields, located below 2 mm of distance from the CGB/GaAs interface and scattered around the CGB. The origin of this structure is not clear but can possibly be attributed to the irregular relaxation of stress during CGB growth. However, as shown in the white circle of Fig. 2(a), a LS filters a TD. Consequently, the authors can apply the idea of TD filtering by LS [18] to understand the mechanism of reduction of the TD in the CGB. The distortion of TDs when bent and the resulting reflection of a TD is found in the blue circle of Fig. 2(a). This may also explain the reduction of defects. For a more in depth study of the compositional distribution of InAlSb CGB and thickness of the InSb, a SIMS measurement was taken. As shown in the inset of Fig. 2(b), In and Al atom counts were measured as a function of Cs+ ion sputtering time. The sputtering time when In or Al atom counts approach 0/s was considered as the CGB/GaAs interface. The composition x of InxAl1 xSb was calculated from the normalized intensity of In and Al atom counts and is illustrated in Fig. 2(b). The composition of InAlSb appears to be concave rather than linear, although the In and Al cell temperatures were linearly controlled by an MBE process computer software. This is reasonable considering the fact that flux of the source is not propositional to the source cell temperature. The abrupt decrease of x near the sample surface is due to the 20 nm-thick In0.3Al0.7Sb cap layer. The thickness of InSb was estimated as 0.41 mm where normalized Al atom counts are less than 0.01, and the composition x of InxAl1 xSb is estimated as less than 0.01 (see red arrow in Fig. 2(b) and its inset). From this estimation, it is found that the time when the flux of the Al approach 0 A˚/s is 42 min in the second phase, and the Al cell temperature for no flux out of the cell is calculated as 940 8C, which is in agreement with the measured Al cell temperature for no flux (950 8C) considering errors due to background noise in the flux sensor. Although this can be seen as an indirect support for SIMS measurements, an XRD measurement of the sample and an XRD simulation of the compositional information from the SIMS were compared for a more direct proof of the accuracy of the compositional information measured by the SIMS. As shown in Fig. 2(c), the measured rocking curve is in good agreement with the simulated rocking curves. With the SIMS (Fig. 2(b)) and XRD (Fig. 2(c)) information, the thickness of InSb can be identified as illustrated in Fig. 2(a). At this point, it is noteworthy that crystal defects disappear within InAlSb CGB and the structural quality of InSb is very high without the defects in Fig. 2(a). Meanwhile, dislocations are found in the crosssectional TEM image of the lower layer of InSb on SGB [15]. Fig. 3(a) shows the morphological information of 0.41 mm-thick InSb on InAlSb CGB with a 20 nm-thick InAlSb capping layer measured by AFM. The root-mean-square surface (RMS) roughness from a 10 mm 10 mm AFM image is 5.3 nm and the peak wavelength of a 2-dimensional (2-D) power spectral density (PSD) analysis is 5 mm. Compared with those of the SGB [10], the RMS roughness is similar to that of the 2 mm-thick InSb on a 500 nmthick InAlSb SGB/1 mm-thick AlSb/GaAs, and the peak wavelength is similar to that of the 2 mm-thick InSb on no InAlSb SGB/1 mmthick AlSb/GaAs. Considering that a larger mosaicity was found in thicker SGB with the larger RMS roughness and the larger peak wavelength of 2D-PSD [15] and that the peak wavelength of PSD is enlarged with increasing thickness of InSb [20], a smaller mosaicity may be expected in CGB due to the decreased thickness of InSb. Fig. 3(b) illustrates electron mobility as a function of measurement temperature. The dependence of mobility drops in a sigmoidal shape. These trends of temperature dependences are
2929
Fig. 3. (a) Surface 2-dimensional power spectral density analysis and AFM image (10 mm 10 mm) of the sample (inset). (b) Measured electron mobility as a function of measurement temperature and comparison of electron mobility at 300 K with references (inset).
in general accordance with the results for undoped InSb layers grown on GaAs [10,20]. The apparent rapid-drop of electron mobility with decreasing temperature in the low temperature range (80–270 K) can be interpreted as an artifact due to the presence of holes unintentionally caused by a p-type background [20]. The measured m300 K is 46,300 cm2/Vs. Finally, the m300 K of 0.41 mm-thick InSb on InAlSb CGB/GaAs is compared with various thicknesses of InSb on GaAs in the inset of Fig. 3(b) [11,13,15]. As shown there, the m300 K of InSb is explicitly enhanced, even under the assumption that the thickness of InSb is enlarged to 0.6 mm by 50% of error during the estimation of InSb thickness. This is attributed to defect reduction by InAlSb CGB. It should be pointed out that for increased high-speed performance, the thickness of InSb active film should be reduced with increased quality (m300 K). 4. Conclusions An InAlSb CGB was introduced to accommodate a lattice mismatch between GaAs (1 0 0) substrate and InSb. The defects caused by the lattice mismatch were effectively reduced at 2 mm from the interfaces between GaAs substrate and InAlSb CGB. This reduction of defect is attributed to LSs parallel to the GaAs interface which reflect, bend, stop and, as result, filter the TDs. The origin of this structure is not clear, however, can possibly be attributed to the irregular relaxation of stress during CGB growth. The thickness of pure InSb layers was found to be 0.41 mm, which was evaluated from SIMS and XRD measurements. The m300 K of 0.41 mm-thick InSb on top of InAlSb CGB is 46,300 cm2/Vs, which surpasses the m300 K of state-of-the-art InSb grown by other methods with
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similar thickness. We expect that CGB will be useful for high-speed device applications using InSb 2DEGs. Acknowledgments This work was supported mainly by KIST Institutional Programs including Dream Project and partially by NRF grant funded by MEST (No. 2011-0016471). References [1] A. Oral, et al. IEEE Trans. Magn. 38 (2002) 2438. [2] S.A. Solin, et al. Appl. Phys. Lett. 80 (2002) 4012. [3] S.A. Solin, et al. Science 289 (2000) 1530.
[4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20]
E. Michel, et al. IEEE Photonics Technol. Lett. 8 (1996) 673. W. Dobbelaere, et al. Appl. Phys. Lett. 60 (1992) 29. S. Ozer, C. Besikci, J. Phys. D: Appl. Phys. 36 (2003) 559. B. Bennett, et al. Solid State Electron. 49 (2005) 1875. A. Joullie´, et al. Mater. Res. Bull. 9 (1974) 241. Z. Burshtein, M. Azulay, Mater. Res. Bull. 19 (1984) 49. J. So¨derstro¨m, et al. Semicond. Sci. Technol. 7 (1992) 337. E. Michel, et al. Appl. Phys. Lett. 65 (1994) 3338. F.E. Ejeckam, et al. Appl. Phys. Lett. 71 (1997) 776. T. Zhang, et al. Appl. Phys. Lett. 84 (2004) 4463. M.C. Debnath, et al. J. Cryst. Growth 267 (2004) 17. T. Sato, et al. Physica E 21 (2004) 615. A. Okamoto, et al. J. Cryst. Growth 227 (2001) 619. T.D. Mishima, et al. Physica B 376 (2006) 591. T.D. Mishima, et al. Appl. Phys. Lett. 88 (2006) 191908. J.Y. Lim, et al. J. Korean Phys. Soc. 55 (2009) 1525. A. Okamoto, I. Shibasaki, J. Cryst. Growth 251 (2003) 560.