**III-V Epitaxy**

[14] Morishita J, Wu S, Ishihara Y, Kimijima T. Observation of purifier performance to reduce moisture in hydrogen chloride by near infrared laser absorption spectrometry. Journal of

[15] Gibian TG, McKinney DS. Infrared spectra and force constants of chloroform and trichlorosilane. Journal of the American Chemical Society. 1951;73:1431-1434

[16] Ushio Electric https://www.ushio.co.jp/jp/feature/halogenHeater/ss05.html (access date:

[17] Habuka H, Nagoya T, Katayama M, Shimada M, Okuyama K. Modelling of epitaxial silicon thin-film growth on a rotating substrate in a horizontal single-wafer reactor.

[19] Haynes WM. CRC Handbook of Chemistry and Physics. 92nd ed. Boca Raton, USA: CRC

[20] Habuka H, Suzuki T, Sakurai T, Negishi Y, Takeuchi T. Non-empirical design of rapid thermal processing system. Japanese Journal of Applied Physics. 2001;40(12):7123-7128

[18] Crippa D, Rode DL, Masi M. Silicon Epitaxy. San Diego, USA: Academic Press; 2001

Journal of Electrochemical Society. 1995;142(12):4272-4278

Applied Physics. 1997;36:L1706

June 23, 2017)

152 Epitaxy

Press; 2012

**Chapter 7**

**Provisional chapter**

**GaN and InN Hexagonal Microdisks**

**GaN and InN Hexagonal Microdisks**

DOI: 10.5772/intechopen.70120

The high-quality GaN microdisks with InGaN/GaN quantum wells (QWs) and InN

taxy (PA-MBE). The samples were analysed using scanning electron microscopy, X-ray diffraction, photoluminescence, cathodoluminescence and high-resolution transmission electron microscope. The characteristics of the GaN microdisks and InN microdisks were

III-nitride materials have been extensively studied for the applications to high-efficiency lighting sources such as light-emitting diodes (LEDs) or spintronics [1–7]. From the changing of

whole visible-light spectrum. **Figure 1** shows the diagram of band-gap energies of III-nitrides with bowing parameters [8] versus lattice constants. However, it is difficult to grow high-

GaN (*a*GaN = 0.3189 nm, *c*GaN = 0.5185 nm) and InN (*a*InN = 0.35446 nm, *c*InN = 0.57034 nm) [8, 9]. Furthermore, because of the high volatility of indium atom at high temperature, it is hard

high-temperature growth techniques (e.g. *T* > 1000°C) such as vapour phase epitaxy (VPE) or

epilayer with an indium concentration higher than 20% is regarded as a high challenge [10]. To overcome these difficulties, a plasma-assisted molecular beam epitaxy (PA-MBE)

metalorganic chemical vapour deposition (MOCVD). The growth of high-quality In*<sup>x</sup>*

Ga1−*<sup>x</sup>*

substrates by plasma-assisted molecular beam epi-

N can be tuned from 0.7 to 3.4 eV to cover the

N/GaN QW thin film by using

Ga1−*<sup>x</sup>* N

Chen-Chi Yang, Ikai Lo, Yu-Chi Hsu and

Chen-Chi Yang, Ikai Lo, Yu-Chi Hsu and

http://dx.doi.org/10.5772/intechopen.70120

microdisks were grown on γ-LiAlO<sup>2</sup>

indium content (*x*), the band-gap of In*<sup>x</sup>*

to grow a homogenous high-indium-concentration In*<sup>x</sup>*

Ga1−*<sup>x</sup>*

studied and the effect of growth temperature was evaluated.

**Keywords:** GaN, InN, microdisk, molecular beam epitaxy

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution,

© 2018 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

N/GaN quantum wells (QWs) because of the large lattice mismatch between

Ga1−*<sup>x</sup>*

N epilayer at lower temperatures, and some substrates

and reproduction in any medium, provided the original work is properly cited.

Hong-Yi Yang

**Abstract**

**1. Introduction**

quality In*<sup>x</sup>*

Ga1−*<sup>x</sup>*

technique was used to grow In*<sup>x</sup>*

Hong-Yi Yang

**Provisional chapter**

## **GaN and InN Hexagonal Microdisks**

**GaN and InN Hexagonal Microdisks**

Chen-Chi Yang, Ikai Lo, Yu-Chi Hsu and Hong-Yi Yang Hong-Yi Yang Additional information is available at the end of the chapter

Chen-Chi Yang, Ikai Lo, Yu-Chi Hsu and

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/intechopen.70120

#### **Abstract**

The high-quality GaN microdisks with InGaN/GaN quantum wells (QWs) and InN microdisks were grown on γ-LiAlO<sup>2</sup> substrates by plasma-assisted molecular beam epitaxy (PA-MBE). The samples were analysed using scanning electron microscopy, X-ray diffraction, photoluminescence, cathodoluminescence and high-resolution transmission electron microscope. The characteristics of the GaN microdisks and InN microdisks were studied and the effect of growth temperature was evaluated.

DOI: 10.5772/intechopen.70120

**Keywords:** GaN, InN, microdisk, molecular beam epitaxy

## **1. Introduction**

III-nitride materials have been extensively studied for the applications to high-efficiency lighting sources such as light-emitting diodes (LEDs) or spintronics [1–7]. From the changing of indium content (*x*), the band-gap of In*<sup>x</sup>* Ga1−*<sup>x</sup>* N can be tuned from 0.7 to 3.4 eV to cover the whole visible-light spectrum. **Figure 1** shows the diagram of band-gap energies of III-nitrides with bowing parameters [8] versus lattice constants. However, it is difficult to grow highquality In*<sup>x</sup>* Ga1−*<sup>x</sup>* N/GaN quantum wells (QWs) because of the large lattice mismatch between GaN (*a*GaN = 0.3189 nm, *c*GaN = 0.5185 nm) and InN (*a*InN = 0.35446 nm, *c*InN = 0.57034 nm) [8, 9]. Furthermore, because of the high volatility of indium atom at high temperature, it is hard to grow a homogenous high-indium-concentration In*<sup>x</sup>* Ga1−*<sup>x</sup>* N/GaN QW thin film by using high-temperature growth techniques (e.g. *T* > 1000°C) such as vapour phase epitaxy (VPE) or metalorganic chemical vapour deposition (MOCVD). The growth of high-quality In*<sup>x</sup>* Ga1−*<sup>x</sup>* N epilayer with an indium concentration higher than 20% is regarded as a high challenge [10]. To overcome these difficulties, a plasma-assisted molecular beam epitaxy (PA-MBE) technique was used to grow In*<sup>x</sup>* Ga1−*<sup>x</sup>* N epilayer at lower temperatures, and some substrates

Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2018 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons

technique. Before mounting on a holder, the LAO substrate was cleaned with acetone (5 min), isopropanol (5 min) and de-ionized water for a while, and then dried with nitrogen gas immediately. After the chemical cleaning, a thermal treatment was introduced to the LAO substrate in the MBE chamber before epitaxial growth. The LAO substrate was out-gassed at 680°C for 10 min. The temperature was defined by a thermal couple equipped with the backside of the substrate. Thereafter, the substrate temperature was decreased to growth temperatures. The Ga wetting layer was performed on the LAO substrate for 5 min at 630°C, and then the two-step method (i.e. two different N/Ga flux ratios from 28.9 to 139.7, for 35 and 70 min, respectively) was used to fabricate the GaN epi-film at 620°C. The flux ratio was represented by beam equivalent pressure (BEP) of evaporative III-group sources from the standard effu-

developed a back process to fabricate an electrical contact for the GaN hexagonal microdisk on a transparent p-type GaN template [12]. In this chapter, we have consistently grown a sample of GaN microdisks to demonstrate the self-assembling model [5], as shown in **Figure 2**.

The surface morphology of the GaN microdisk sample was evaluated by the field emission scanning electron microscopy (FE-SEM, SII-3050). **Figure 2** shows SEM images with a tilted

**Figure 2.** (a)–(c) The tilt-view SEM images of the sample, the scale bars are 10, 4 and 1 μm, respectively. (d)–(f) The top-

view SEM images of the sample, the scale bars are 10, 4 and 1 μm, respectively.

source from the rf-plasma cell [11]. In our previous study, we

1) hexagonal microdisks [5]. Besides, we

GaN and InN Hexagonal Microdisks http://dx.doi.org/10.5772/intechopen.70120 157

sion cell against that of the N<sup>2</sup>

**2.2. Characteristics of GaN microdisks**

showed the characteristics of *c*-plane GaN (<sup>0</sup> <sup>0</sup> <sup>0</sup>¯

**Figure 1.** The diagram shows band-gap energies of III-nitrides with bowing parameters versus lattice constants. The substrate materials were also presented.

(such as LiAlO<sup>2</sup> , LiGaO<sup>2</sup> , ZnO) were selected to grow GaN epilayers in order to minimize the lattice mismatch between the substrate and GaN layer as compared to the commercial substrates, e.g. sapphire, SiC or Si (1 1 1). From the values of wurtzite GaN on JCPDS files No. 50-0792 and those of γ-LiAlO<sup>2</sup> on No. 38-1464, the lattice mismatch between *c*/2LAO and *a*GaN is 1.5% indicating that it is suit to grow *c*-plane GaN on the g-LiAlO<sup>2</sup> substrate. In this chapter, we show the high-quality epitaxial growth of GaN microdisks with In*<sup>x</sup>* Ga1−*<sup>x</sup>* N/GaN quantum wells and InN microdisks on γ-LiAlO<sup>2</sup> lithium aluminium oxide (LAO) substrates at low temperatures by the plasma-assisted molecular beam epitaxy system. Consequently, GaN and InN microdisks provide better opportunities to fabricate the In*<sup>x</sup>* Ga1−*<sup>x</sup>* N/GaN microdisk quantum wells for the application of full-colour micron LED without the sapphire substrate, which is mostly used for the bulk GaN-based quantum wells in commercial LEDs but has a larger lattice mismatch with InN [1, 2].

#### **2. GaN hexagonal microdisks**

#### **2.1. Growth of GaN hexagonal microdisks**

The sample was grown on a high-quality 1 **×** 1 cm2 LAO (1 0 0) substrate by using a lowtemperature PA-MBE system (Veeco Applied-GEN 930) with standard effusion cells for Ga-evaporation and an rf-plasma cell with 450 W for N<sup>2</sup> -plasma source. The LAO substrate was cut from the crystal ingot, which was fabricated by the traditional Czochralski pulling technique. Before mounting on a holder, the LAO substrate was cleaned with acetone (5 min), isopropanol (5 min) and de-ionized water for a while, and then dried with nitrogen gas immediately. After the chemical cleaning, a thermal treatment was introduced to the LAO substrate in the MBE chamber before epitaxial growth. The LAO substrate was out-gassed at 680°C for 10 min. The temperature was defined by a thermal couple equipped with the backside of the substrate. Thereafter, the substrate temperature was decreased to growth temperatures. The Ga wetting layer was performed on the LAO substrate for 5 min at 630°C, and then the two-step method (i.e. two different N/Ga flux ratios from 28.9 to 139.7, for 35 and 70 min, respectively) was used to fabricate the GaN epi-film at 620°C. The flux ratio was represented by beam equivalent pressure (BEP) of evaporative III-group sources from the standard effusion cell against that of the N<sup>2</sup> source from the rf-plasma cell [11]. In our previous study, we showed the characteristics of *c*-plane GaN (<sup>0</sup> <sup>0</sup> <sup>0</sup>¯ 1) hexagonal microdisks [5]. Besides, we developed a back process to fabricate an electrical contact for the GaN hexagonal microdisk on a transparent p-type GaN template [12]. In this chapter, we have consistently grown a sample of GaN microdisks to demonstrate the self-assembling model [5], as shown in **Figure 2**.

#### **2.2. Characteristics of GaN microdisks**

(such as LiAlO<sup>2</sup>

156 Epitaxy

, LiGaO<sup>2</sup>

No. 50-0792 and those of γ-LiAlO<sup>2</sup>

substrate materials were also presented.

larger lattice mismatch with InN [1, 2].

**2. GaN hexagonal microdisks**

**2.1. Growth of GaN hexagonal microdisks**

The sample was grown on a high-quality 1 **×** 1 cm2

Ga-evaporation and an rf-plasma cell with 450 W for N<sup>2</sup>

, ZnO) were selected to grow GaN epilayers in order to minimize

on No. 38-1464, the lattice mismatch between *c*/2LAO and

Ga1−*<sup>x</sup>*

LAO (1 0 0) substrate by using a low-


substrate. In this

N/GaN

Ga1−*<sup>x</sup>*

N/GaN microdisk

the lattice mismatch between the substrate and GaN layer as compared to the commercial substrates, e.g. sapphire, SiC or Si (1 1 1). From the values of wurtzite GaN on JCPDS files

**Figure 1.** The diagram shows band-gap energies of III-nitrides with bowing parameters versus lattice constants. The

quantum wells and InN microdisks on γ-LiAlO<sup>2</sup> lithium aluminium oxide (LAO) substrates at low temperatures by the plasma-assisted molecular beam epitaxy system. Consequently, GaN

quantum wells for the application of full-colour micron LED without the sapphire substrate, which is mostly used for the bulk GaN-based quantum wells in commercial LEDs but has a

temperature PA-MBE system (Veeco Applied-GEN 930) with standard effusion cells for

was cut from the crystal ingot, which was fabricated by the traditional Czochralski pulling

*a*GaN is 1.5% indicating that it is suit to grow *c*-plane GaN on the g-LiAlO<sup>2</sup>

and InN microdisks provide better opportunities to fabricate the In*<sup>x</sup>*

chapter, we show the high-quality epitaxial growth of GaN microdisks with In*<sup>x</sup>*

The surface morphology of the GaN microdisk sample was evaluated by the field emission scanning electron microscopy (FE-SEM, SII-3050). **Figure 2** shows SEM images with a tilted

**Figure 2.** (a)–(c) The tilt-view SEM images of the sample, the scale bars are 10, 4 and 1 μm, respectively. (d)–(f) The topview SEM images of the sample, the scale bars are 10, 4 and 1 μm, respectively.

angle and a top-view of the sample, respectively. The morphology of the sample exhibited that a two-dimensional (2D) *M*-plane GaN film and three-dimensional (3D) *c*-plane GaN hexagonal microdisks were grown on the LAO substrate. The micrographic images of the sample showed that the 2D *M*-plane GaN epi-film was developed along with the lateral orientation [<sup>1</sup> <sup>1</sup> ¯ 2 0]GaN // [0 0 1]LAO, while the 3D *c*-plane GaN hexagonal microdisks were grown atop an anionic hexagonal basal plane of LAO. The two-orientation growth of GaN nanopillars on the LAO substrate has been reported in our previous papers [13, 14]. **Figure 2(c)** shows that the neck of contact area between the GaN microdisk and the LAO substrate is small (e.g. less than 200 nm). In addition, the lattice mismatch between *c*/2LAO and *a*GaN is only 1.5%. It implies that the GaN microdisk is nearly freestanding as a new substrate for further growth of In*<sup>x</sup>* Ga1−*<sup>x</sup>* N/GaN quantum well (QW) on the top, which ignores the lattice mismatch against the LAO substrate. **Figure 2(f)** shows the enlarged SEM image with a top-view of the GaN hexagonal microdisks shown in the centre of **Figure 2(d)**, and the diameter of the GaN microdisk is about 2.0 μm. Based on the self-assembling model, we will extend to the growth of In*x* Ga1−*<sup>x</sup>* N/GaN QW on GaN microdisk and show its characteristics in Section 4.

The optical properties of the GaN microdisk sample were measured by photoluminescence spectroscopy (PL, HORIBA HR800) at room temperature with a light source of He-Cd 325 nm laser. We performed the laser beam focusing on two different spots (S1 and S2) and compared the results with the spot without any microdisk (i.e. mostly *M*-plane GaN and labelled as background), as shown in **Figure 3**. Two major peaks were obtained for each measurement (S1 or S2). These two major peaks were confirmed by a non-linear Gaussian-function curve fitting with the software Origin (Pro. 8.0). The result of the non-linear Gaussian-function curve fitting showed that the positions of two major peaks for two spots were very consistent. The averaged value for the first peak is (3.385 ± 0.001) eV with the full width at half maximum (FWHM) value equal to (0.128 ± 0.001) eV. It is due to the band-gap transition of wurtzite GaN. The averaged value for the second peak is (2.226 ± 0.003) eV with the FWHM value equal to (0.363 ± 0.033) eV. It is an energy level related to structural defects (e.g. YL in reference [15]) in GaN, so that the FWHM value of GaN is smaller than that of the defect level. The PL intensity corresponding to wurtzite GaN and the defect level indicates that *c*-plane GaN microdisk is a higher quality structure than *M*-plane GaN because the intensity of defect level from GaN microdisk is lower than that from *M*-plane GaN background.

between the {<sup>0</sup> <sup>0</sup> <sup>0</sup>¯

plane background of the sample.

of [<sup>1</sup> <sup>1</sup> <sup>2</sup>¯ <sup>0</sup>]

of Ga atoms by ¯¯

sive width along the [<sup>1</sup> <sup>1</sup>¯ <sup>0</sup> <sup>0</sup>]

the (1¯

and the *d*-spacing between the {1¯

file No. 50-0792) with *a* = *b* = 0.3189, *c* = 0.5185 and *u* = ¯¯

1} planes of GaN hexagonal microdisk was measured to be *d***<sup>c</sup>**

GaN // [0 0 1]LAO were performed at the areas HR01 and HR02 of the sample, as shown

*α*-bonds of N atoms to form the microdisk [5]. The lateral over-growth along

GaN direction per unit step-layer was equal to <sup>√</sup>

*d***M** = 0.2785 nm. Compared with the values on JCPDS file No. 50-0792 which are 0.5185 and 0.2762 nm, respectively, the difference between wurtzite GaN microdisk and bulk GaN for *d***<sup>c</sup>** and *d***M** are 1.33 and 0.08%, respectively, revealing that the lattice constant of GaN microdisk is slightly larger than that of bulk GaN. The high-resolution TEM images with the beam direction

**Figure 3.** The PL spectra measured at room temperature for different spots (S1 and S2) of the GaN microdisks and *M*-

in **Figure 4(a)**. The symmetric hexagonal shape reveals the high-quality crystalline structure of the GaN microdisk, as shown in **Figure 4(c)**. The angle between the edge and the growth direction can be examined directly by the high-resolution TEM image performed at HR02 to be about 28°, as shown in **Figure 4(d)**. The ball-stick model for the standard wurtzite GaN (JCPDS

GaN hexagonal microdisk in **Figure 4(e)**, where blue balls represented Ga atoms and red balls represented N atoms. The *c*-plane GaN (<sup>0</sup> <sup>0</sup> <sup>0</sup> <sup>1</sup>¯) hexagonal microdisk was built up with the capture of N atoms by the ¯¯*β*-dangling bonds of the most-outside Ga atoms and then the capture

1 0 0) direction was extended to one *d***M**-spacing for each unit step-layer (i.e. *d***<sup>c</sup>**

resulting in the angle of 28° off the *c*-axis. Based on the ball-stick model, the laterally exten-

1 0 0} planes of GaN hexagonal microdisk was found to be

*α*/*c* = 3/8 was used to simulate the *c*-plane

\_\_ \_\_3 = 0.5254 nm

GaN and InN Hexagonal Microdisks http://dx.doi.org/10.5772/intechopen.70120 159


<sup>2</sup> *<sup>a</sup>*. The edge was

The microstructure of the GaN microdisk sample was analysed by field emission transmission electron microscopy (FE-TEM) (Phillips, model Tecnai F-20) with an electron voltage of 200 kV. The cross-sectional TEM specimen of the sample was prepared by a dual-beam Focus Ion Beam system (FIB, Seiko Inc., SII-3050), on the cleavage plane along the [<sup>1</sup> <sup>1</sup>¯ <sup>0</sup> <sup>0</sup>] direction of the *c*-plane GaN hexagonal microdisk. The FIB was performed with an accelerated voltage of 30 kV to cut the samples roughly and then refined the specimen further by an accelerated voltage of 5 kV. **Figure 4(a)** shows the bright field image with [<sup>1</sup> <sup>1</sup> ¯ 2 0] GaN // [0 0 1]LAO zone axis. It clearly exhibited that the GaN microdisk was well formed on the LAO substrate. The height for the *c*-plane GaN hexagonal microdisk from neck to top was about 4.1 μm. The selective area diffraction (SAD) pattern at the top area of the GaN microdisk shown in **Figure 4(b)** clearly showed one single rectangular diffraction pattern at the location of DP01, indicating that the hexagonal microdisk was uniquely formed by the *c*-plane wurtzite GaN. The *d*-spacing

angle and a top-view of the sample, respectively. The morphology of the sample exhibited that a two-dimensional (2D) *M*-plane GaN film and three-dimensional (3D) *c*-plane GaN hexagonal microdisks were grown on the LAO substrate. The micrographic images of the sample showed that the 2D *M*-plane GaN epi-film was developed along with the lateral orientation

2 0]GaN // [0 0 1]LAO, while the 3D *c*-plane GaN hexagonal microdisks were grown atop an anionic hexagonal basal plane of LAO. The two-orientation growth of GaN nanopillars on the LAO substrate has been reported in our previous papers [13, 14]. **Figure 2(c)** shows that the neck of contact area between the GaN microdisk and the LAO substrate is small (e.g. less than 200 nm). In addition, the lattice mismatch between *c*/2LAO and *a*GaN is only 1.5%. It implies that the GaN microdisk is nearly freestanding as a new substrate for further growth

N/GaN quantum well (QW) on the top, which ignores the lattice mismatch against

the LAO substrate. **Figure 2(f)** shows the enlarged SEM image with a top-view of the GaN hexagonal microdisks shown in the centre of **Figure 2(d)**, and the diameter of the GaN microdisk is about 2.0 μm. Based on the self-assembling model, we will extend to the growth of

The optical properties of the GaN microdisk sample were measured by photoluminescence spectroscopy (PL, HORIBA HR800) at room temperature with a light source of He-Cd 325 nm laser. We performed the laser beam focusing on two different spots (S1 and S2) and compared the results with the spot without any microdisk (i.e. mostly *M*-plane GaN and labelled as background), as shown in **Figure 3**. Two major peaks were obtained for each measurement (S1 or S2). These two major peaks were confirmed by a non-linear Gaussian-function curve fitting with the software Origin (Pro. 8.0). The result of the non-linear Gaussian-function curve fitting showed that the positions of two major peaks for two spots were very consistent. The averaged value for the first peak is (3.385 ± 0.001) eV with the full width at half maximum (FWHM) value equal to (0.128 ± 0.001) eV. It is due to the band-gap transition of wurtzite GaN. The averaged value for the second peak is (2.226 ± 0.003) eV with the FWHM value equal to (0.363 ± 0.033) eV. It is an energy level related to structural defects (e.g. YL in reference [15]) in GaN, so that the FWHM value of GaN is smaller than that of the defect level. The PL intensity corresponding to wurtzite GaN and the defect level indicates that *c*-plane GaN microdisk is a higher quality structure than *M*-plane GaN because the intensity of defect level from GaN microdisk is lower

The microstructure of the GaN microdisk sample was analysed by field emission transmission electron microscopy (FE-TEM) (Phillips, model Tecnai F-20) with an electron voltage of 200 kV. The cross-sectional TEM specimen of the sample was prepared by a dual-beam Focus Ion Beam system (FIB, Seiko Inc., SII-3050), on the cleavage plane along the [<sup>1</sup> <sup>1</sup>¯ <sup>0</sup> <sup>0</sup>] direction of the *c*-plane GaN hexagonal microdisk. The FIB was performed with an accelerated voltage of 30 kV to cut the samples roughly and then refined the specimen further by an accelerated

axis. It clearly exhibited that the GaN microdisk was well formed on the LAO substrate. The height for the *c*-plane GaN hexagonal microdisk from neck to top was about 4.1 μm. The selective area diffraction (SAD) pattern at the top area of the GaN microdisk shown in **Figure 4(b)** clearly showed one single rectangular diffraction pattern at the location of DP01, indicating that the hexagonal microdisk was uniquely formed by the *c*-plane wurtzite GaN. The *d*-spacing

2 0]

GaN // [0 0 1]LAO zone

voltage of 5 kV. **Figure 4(a)** shows the bright field image with [<sup>1</sup> <sup>1</sup> ¯

N/GaN QW on GaN microdisk and show its characteristics in Section 4.

[<sup>1</sup> <sup>1</sup> ¯

158 Epitaxy

of In*<sup>x</sup>*

In*x* Ga1−*<sup>x</sup>*

Ga1−*<sup>x</sup>*

than that from *M*-plane GaN background.

**Figure 3.** The PL spectra measured at room temperature for different spots (S1 and S2) of the GaN microdisks and *M*plane background of the sample.

between the {<sup>0</sup> <sup>0</sup> <sup>0</sup>¯ 1} planes of GaN hexagonal microdisk was measured to be *d***<sup>c</sup>** = 0.5254 nm and the *d*-spacing between the {1¯ 1 0 0} planes of GaN hexagonal microdisk was found to be *d***M** = 0.2785 nm. Compared with the values on JCPDS file No. 50-0792 which are 0.5185 and 0.2762 nm, respectively, the difference between wurtzite GaN microdisk and bulk GaN for *d***<sup>c</sup>** and *d***M** are 1.33 and 0.08%, respectively, revealing that the lattice constant of GaN microdisk is slightly larger than that of bulk GaN. The high-resolution TEM images with the beam direction of [<sup>1</sup> <sup>1</sup> <sup>2</sup>¯ <sup>0</sup>] GaN // [0 0 1]LAO were performed at the areas HR01 and HR02 of the sample, as shown in **Figure 4(a)**. The symmetric hexagonal shape reveals the high-quality crystalline structure of the GaN microdisk, as shown in **Figure 4(c)**. The angle between the edge and the growth direction can be examined directly by the high-resolution TEM image performed at HR02 to be about 28°, as shown in **Figure 4(d)**. The ball-stick model for the standard wurtzite GaN (JCPDS file No. 50-0792) with *a* = *b* = 0.3189, *c* = 0.5185 and *u* = ¯¯ *α*/*c* = 3/8 was used to simulate the *c*-plane GaN hexagonal microdisk in **Figure 4(e)**, where blue balls represented Ga atoms and red balls represented N atoms. The *c*-plane GaN (<sup>0</sup> <sup>0</sup> <sup>0</sup> <sup>1</sup>¯) hexagonal microdisk was built up with the capture of N atoms by the ¯¯*β*-dangling bonds of the most-outside Ga atoms and then the capture of Ga atoms by ¯¯ *α*-bonds of N atoms to form the microdisk [5]. The lateral over-growth along the (1¯ 1 0 0) direction was extended to one *d***M**-spacing for each unit step-layer (i.e. *d***<sup>c</sup>** -spacing), resulting in the angle of 28° off the *c*-axis. Based on the ball-stick model, the laterally extensive width along the [<sup>1</sup> <sup>1</sup>¯ <sup>0</sup> <sup>0</sup>] GaN direction per unit step-layer was equal to <sup>√</sup> \_\_ \_\_3 <sup>2</sup> *<sup>a</sup>*. The edge was

substrate with an InGaN buffer layer between them by a low-temperature PA-MBE system (Veeco Applied-GEN 930). The LAO substrate was cut from the crystal ingot, which was fabricated by the traditional Czochralski pulling technique. The growth details were described

The crystal structure of the InN microdisk sample was characterized by the high-resolution X-ray diffraction (XRD; Bede D1) measurement and is shown in **Figure 5(a)**. From the result of X-ray diffraction pattern (i.e. the peak at *2θ* = 31.69°), we estimated the indium content of

diffraction patterns were obtained and matched with those data of the standard wurtzite structure bulk InN (JCPDS file No. 50-1239) by the asymmetric double sigmoidal linear curve fitting

was evaluated to be *d***0002** = 0.28216 nm from the Bragg's law (*2d*sin*θ* = *n*λ) with Cu Kα1wavelength *λ* = 0.1540562 nm. The lattice constant of wurtzite InN microdisk is smaller than that of bulk InN by comparing with the value on JCPDS file, *d***0002** = 0.28528 nm, and the difference between InN

**Figure 5.** (a) The X-ray 2Theta-Omega scan of the sample. (b) The top-view SEM image of InN hexagonal thin disk, the

scale bar is 0.5 μm. (c) The tilt-view SEM image of InN hexagonal thin disk, the scale bar is 0.5 μm.

N on the basis of Vegard's law to be about 20% [17].The peaks at *2θ* = 29.07, 31.31, 32.29 and 34.69° represented the X-ray diffraction patterns from the *M*-plane InN (1 <sup>1</sup>¯<sup>0</sup> 0), *c*-plane InN

1 0 0) and LAO (1 0 0), respectively. These peak positions at the X-ray

2} planes of InN

GaN and InN Hexagonal Microdisks http://dx.doi.org/10.5772/intechopen.70120 161

completely in our previous paper which was published in *AIP Advances* [16].

with the software Quick Graph (Version 2.0). The *d*-spacing between the {<sup>0</sup> <sup>0</sup> <sup>0</sup> ¯

**3.2. Characteristics of InN microdisks**

<sup>2</sup>), *M*-plane GaN (1¯

microdisk and bulk InN is 1.09%.

In*x* Ga1−*<sup>x</sup>*

(<sup>0</sup> <sup>0</sup> <sup>0</sup>¯

**Figure 4.** TEM analyses of the GaN hexagonal microdisk: (a) the bright field image with [<sup>1</sup> <sup>1</sup> ¯ 2 0]GaN // [0 0 1]LAO zone axis. The selective area diffraction patterns taken at the point shown in (a) are presented in (b), the scale bars are 2 (1/nm). The high-resolution TEM images taken at the points shown in (a) are presented in (c) and (d), the scale bars are 2 nm and (e) the ball-stick model for GaN hexagonal microdisk.

then tilted off the *c*-axis [0 <sup>0</sup> <sup>0</sup> <sup>1</sup>¯] direction by the angle, *φ* = tan−1( √ \_\_ \_\_3 <sup>2</sup> *<sup>a</sup>*/*c*) = 28.1°, where <sup>√</sup> \_\_ \_\_3 <sup>2</sup> *<sup>a</sup>* is equal to one *d***M**, as shown in **Figure 4(e)**. We also calculated the angle from the measured SAD data at the GaN hexagonal microdisk in **Figure 4(b)**, and obtained that the *d*-spacing between the {<sup>0</sup> <sup>0</sup> <sup>0</sup>¯ 1} planes was *d***<sup>c</sup>** = 0.5254 nm and the *d*-spacing between the {1 <sup>1</sup>¯ <sup>0</sup> <sup>0</sup>} planes was *d***M** = 0.2785 nm, resulting in *φ* = tan−1(*d***M**/*d***<sup>c</sup>** ) = 28.0°, which was in good agreement with the model predicted. The angle between the edge and the growth direction can be examined directly by the high-resolution TEM image performed at HR02 to be about 28°, as shown in **Figure 4(d)**.

## **3. InN hexagonal microdisks**

#### **3.1. Growth of InN hexagonal microdisks**

The two-orientation growth of GaN nanopillars on the LAO substrate has been reported in our previous papers [13, 14] and reconfirmed in Section 2. In this section, we applied the two-orientation growth mechanism to grow the 2D *M*-plane InN epi-film and 3D *c*-plane InN hexagonal thin disks on the LAO substrate with the InGaN buffer layer at low growth temperature (470°C). The InN microdisk sample was grown on a high-quality 1 × 1 cm<sup>2</sup> LAO (1 0 0) substrate with an InGaN buffer layer between them by a low-temperature PA-MBE system (Veeco Applied-GEN 930). The LAO substrate was cut from the crystal ingot, which was fabricated by the traditional Czochralski pulling technique. The growth details were described completely in our previous paper which was published in *AIP Advances* [16].

#### **3.2. Characteristics of InN microdisks**

then tilted off the *c*-axis [0 <sup>0</sup> <sup>0</sup> <sup>1</sup>¯] direction by the angle, *φ* = tan−1(

are 2 nm and (e) the ball-stick model for GaN hexagonal microdisk.

**Figure 4.** TEM analyses of the GaN hexagonal microdisk: (a) the bright field image with [<sup>1</sup> <sup>1</sup> ¯

the {<sup>0</sup> <sup>0</sup> <sup>0</sup>¯

160 Epitaxy

1} planes was *d***<sup>c</sup>**

**3. InN hexagonal microdisks**

**3.1. Growth of InN hexagonal microdisks**

*d***M** = 0.2785 nm, resulting in *φ* = tan−1(*d***M**/*d***<sup>c</sup>**

equal to one *d***M**, as shown in **Figure 4(e)**. We also calculated the angle from the measured SAD data at the GaN hexagonal microdisk in **Figure 4(b)**, and obtained that the *d*-spacing between

axis. The selective area diffraction patterns taken at the point shown in (a) are presented in (b), the scale bars are 2 (1/nm). The high-resolution TEM images taken at the points shown in (a) are presented in (c) and (d), the scale bars

predicted. The angle between the edge and the growth direction can be examined directly by the high-resolution TEM image performed at HR02 to be about 28°, as shown in **Figure 4(d)**.

The two-orientation growth of GaN nanopillars on the LAO substrate has been reported in our previous papers [13, 14] and reconfirmed in Section 2. In this section, we applied the two-orientation growth mechanism to grow the 2D *M*-plane InN epi-film and 3D *c*-plane InN hexagonal thin disks on the LAO substrate with the InGaN buffer layer at low growth temper-

ature (470°C). The InN microdisk sample was grown on a high-quality 1 × 1 cm<sup>2</sup>

√ \_\_ \_\_3

) = 28.0°, which was in good agreement with the model

= 0.5254 nm and the *d*-spacing between the {1 <sup>1</sup>¯ <sup>0</sup> <sup>0</sup>} planes was

<sup>2</sup> *<sup>a</sup>*/*c*) = 28.1°, where <sup>√</sup>

\_\_ \_\_3 <sup>2</sup> *<sup>a</sup>* is

2 0]GaN // [0 0 1]LAO zone

LAO (1 0 0)

The crystal structure of the InN microdisk sample was characterized by the high-resolution X-ray diffraction (XRD; Bede D1) measurement and is shown in **Figure 5(a)**. From the result of X-ray diffraction pattern (i.e. the peak at *2θ* = 31.69°), we estimated the indium content of In*x* Ga1−*<sup>x</sup>* N on the basis of Vegard's law to be about 20% [17].The peaks at *2θ* = 29.07, 31.31, 32.29 and 34.69° represented the X-ray diffraction patterns from the *M*-plane InN (1 <sup>1</sup>¯<sup>0</sup> 0), *c*-plane InN (<sup>0</sup> <sup>0</sup> <sup>0</sup>¯ <sup>2</sup>), *M*-plane GaN (1¯ 1 0 0) and LAO (1 0 0), respectively. These peak positions at the X-ray diffraction patterns were obtained and matched with those data of the standard wurtzite structure bulk InN (JCPDS file No. 50-1239) by the asymmetric double sigmoidal linear curve fitting with the software Quick Graph (Version 2.0). The *d*-spacing between the {<sup>0</sup> <sup>0</sup> <sup>0</sup> ¯ 2} planes of InN was evaluated to be *d***0002** = 0.28216 nm from the Bragg's law (*2d*sin*θ* = *n*λ) with Cu Kα1wavelength *λ* = 0.1540562 nm. The lattice constant of wurtzite InN microdisk is smaller than that of bulk InN by comparing with the value on JCPDS file, *d***0002** = 0.28528 nm, and the difference between InN microdisk and bulk InN is 1.09%.

**Figure 5.** (a) The X-ray 2Theta-Omega scan of the sample. (b) The top-view SEM image of InN hexagonal thin disk, the scale bar is 0.5 μm. (c) The tilt-view SEM image of InN hexagonal thin disk, the scale bar is 0.5 μm.

The surface morphology of the InN microdisk sample was evaluated by the field emission scanning electron microscopy (FE-SEM, SII-3050). **Figure 5(b)** shows the top-view SEM image of the sample, and the diameter of the InN microdisk was about 0.96 μm. The morphology of the sample exhibited that 2D *M*-plane InN epi-film and 3D *c*-plane InN hexagonal microdisks were grown on the LAO substrate. **Figure 5(c)** shows the tilt-view SEM image of the InN microdisk. The micrographic images of the sample showed that the 3D *c*-plane InN hexagonal microdisks and nanopillars were grown atop an anionic hexagonal basal plane of LAO, while the 2D *<sup>M</sup>*-plane InN epi-film was developed along with the lateral orientation [<sup>1</sup> <sup>1</sup> ¯ 2 0] InN // [0 0 1]LAO. Compared with the shape of GaN microdisk, InN microdisk was thinner than GaN microdisk.

The microstructure of the sample was analysed by field emission transmission electron microscopy (FE-TEM) (Phillips, model Tecnai F-20) with an electron voltage of 200 kV. The cross-sectional TEM specimen of the sample was prepared by a dual-beam FIB system (Seiko Inc., SII-3050), on the cleavage plane along the [1¯ 1 0 0] direction of the *c*-plane InN hexagonal thin disk. The FIB was performed with an accelerated voltage of 30 kV to cut the samples roughly and then refined the specimen further by an accelerated voltage of 5 kV. **Figure 6(a)** shows the bright field image with the [<sup>1</sup> <sup>1</sup>¯ 2 0] InN // [0 0 1]LAO zone axis. It clearly exhibited that InN was well-formed on the InGaN buffer layer, and the InGaN buffer layer was wellestablished on the GaN epi-layer. The thicknesses of *M*-plane InN, *M*-plane InGaN and *M*plane GaN were measured to be about 265, 51 and 137 nm, respectively. The height for the *c*-plane InN hexagonal thin disk from neck to top was about 188 nm. The *c*-plane wurtzite structure was followed up to the neck area and formed a uniform *c*-plane InGaN pyramidshaped structure. Outside the pyramid-shaped structure, the wave-shaped structures were produced by the staking faults between the misfit *c*-plane wurtzite structures of InGaN and InN. The wave-shaped structures became uniform and then the *c*-plane wurtzite structure was followed further to form the InN hexagonal microdisk. The selective area diffraction (SAD) pattern at the top area of hexagonal thin disk shown in **Figure 6(b)** clearly showed one single rectangular diffraction pattern at the location of DP01, indicating that the hexagonal thin disk was uniquely formed by the *c*-plane wurtzite InN. The *d*-spacing between the {<sup>0</sup> <sup>0</sup> <sup>0</sup> ¯ 1} planes of InN hexagonal thin disk was measured to be *d***<sup>c</sup>** = 0.5687 nm and the *d*-spacing between the {1¯ 1 0 0} planes of InN hexagonal thin disk was *d***M** = 0.3025 nm. Compared with the values on JCPDS file No. 50-1239, which are 0.5703 and 0.30647 nm, respectively, the difference between wurtzite InN microdisk and bulk InN for *d***<sup>c</sup>** and *d***M** are 0.28 and 1.24%, respectively, revealing that the lattice constant of wurtzite InN microdisk is smaller than that of bulk InN. The angle between edge and growth direction can be examined directly by the high-resolution TEM image performed at HR01 to be about 73°, as shown in **Figure 6(c)**. To establish the growth mechanism of the thin InN hexagonal microdisk, we demonstrated a ball-stick model for the self-assembled thin InN microdisk. The ball-stick model for the standard wurtzite InN (JCPDS file No. 50-1239) with *a* = *b* = 0.3537, *c* = 0.5703 and *<sup>u</sup>* <sup>=</sup> ¯¯ *α*/*c* = 3/8 was used to simulate the *c*-plane InN hexagonal microdisk, as shown in **Figure 6(d)** and **(e)**, where blue balls represented In atoms and red balls represented N atoms. In the case of InN thin disk, when the growth temperature lowered to 470°C, the *c*-plane InN (<sup>0</sup> <sup>0</sup> <sup>0</sup> ¯ 1) hexagonal thin disk was built up with the capture of N atoms by the ¯¯*β*-dangling bonds of the most-outside In atoms and then the lateral over-growth occurred; and the capture of In atoms by ¯¯*β*-dangling bonds of N atoms to form

the thin microdisk. The lateral over-growth along the (1¯

\_\_

\_\_

**Figure 6(b)**, and obtained that the *d*-spacing between the {<sup>0</sup> <sup>0</sup> <sup>0</sup>¯

Based on the ball-stick model, the laterally extensive width along the [1¯

**Figure 6.** TEM analyses of the InN hexagonal thin disk: (a) the bright field image with [<sup>1</sup> <sup>1</sup> ¯

3 *a*/*c*) = 72.76°, where 3 √

) = 72.60°, which was in good agreement with the model predicted (**Table 1**).

*d***M**-spacings for each unit step-layer (i.e. *d***<sup>c</sup>**

model for InN epilayer: (d) the chemical bonds of the (0 <sup>0</sup> <sup>0</sup>¯

unit step-layer was equal to 3 √

and the *d*-spacing between the {1¯

**4. InGaN/GaN quantum well**

by the angle, *φ* = tan−1(3 √

*d***c**

**4.1. Growth**

1 0 0) direction was extended to six

1 0 0]

<sup>3</sup> *<sup>a</sup>* is equal to 6*d***M**, as shown in **Figure 6(e)**.

1} planes was *d***<sup>c</sup>**

InN direction per

2 0 1]InN // [0 0 1]LAO zone axis.

GaN and InN Hexagonal Microdisks http://dx.doi.org/10.5772/intechopen.70120 163

1] direction

= 0.5687 nm


1 0 0} planes was *d***<sup>M</sup>** = 0.3025 nm, resulting in *φ* = tan−1(6*d***M**/

1) surface and (e) the hexagonal thin disk.

<sup>3</sup> *<sup>a</sup>*. The edge was then tilted off the *c*-axis [0 <sup>0</sup> <sup>0</sup>¯

\_\_

We also calculated the angle from the measured SAD data at the InN hexagonal thin disk in

The selective area diffraction patterns taken at the points shown in (a) are presented in (b), the scale bars are 2 (1/nm). The high-resolution TEM images taken at the point shown in (a) are presented in (c), the scale bars are 2 nm. The ball-stick

The growth mechanism of the awl-shaped GaN microdisk is divergently self-assembled, indicating that the hexagonal neck area for initial nucleation between GaN microdisk and LAO substrate is very small (diameter ~100 nm), and the strain due to the lattice-mismatch between GaN and LAO substrate will not be delivered to the awl-shaped GaN microdisk at the top. This is the way that the GaN microdisk can be grown in balance with a good awl-shape of

**Figure 6.** TEM analyses of the InN hexagonal thin disk: (a) the bright field image with [<sup>1</sup> <sup>1</sup> ¯ 2 0 1]InN // [0 0 1]LAO zone axis. The selective area diffraction patterns taken at the points shown in (a) are presented in (b), the scale bars are 2 (1/nm). The high-resolution TEM images taken at the point shown in (a) are presented in (c), the scale bars are 2 nm. The ball-stick model for InN epilayer: (d) the chemical bonds of the (0 <sup>0</sup> <sup>0</sup>¯ 1) surface and (e) the hexagonal thin disk.

the thin microdisk. The lateral over-growth along the (1¯ 1 0 0) direction was extended to six *d***M**-spacings for each unit step-layer (i.e. *d***<sup>c</sup>** -spacing), resulting in the angle of 73° off the *c*-axis. Based on the ball-stick model, the laterally extensive width along the [1¯ 1 0 0] InN direction per unit step-layer was equal to 3 √ \_\_ <sup>3</sup> *<sup>a</sup>*. The edge was then tilted off the *c*-axis [0 <sup>0</sup> <sup>0</sup>¯ 1] direction by the angle, *φ* = tan−1(3 √ \_\_ 3 *a*/*c*) = 72.76°, where 3 √ \_\_ <sup>3</sup> *<sup>a</sup>* is equal to 6*d***M**, as shown in **Figure 6(e)**. We also calculated the angle from the measured SAD data at the InN hexagonal thin disk in **Figure 6(b)**, and obtained that the *d*-spacing between the {<sup>0</sup> <sup>0</sup> <sup>0</sup>¯ 1} planes was *d***<sup>c</sup>** = 0.5687 nm and the *d*-spacing between the {1¯ 1 0 0} planes was *d***<sup>M</sup>** = 0.3025 nm, resulting in *φ* = tan−1(6*d***M**/ *d***c** ) = 72.60°, which was in good agreement with the model predicted (**Table 1**).

## **4. InGaN/GaN quantum well**

#### **4.1. Growth**

The surface morphology of the InN microdisk sample was evaluated by the field emission scanning electron microscopy (FE-SEM, SII-3050). **Figure 5(b)** shows the top-view SEM image of the sample, and the diameter of the InN microdisk was about 0.96 μm. The morphology of the sample exhibited that 2D *M*-plane InN epi-film and 3D *c*-plane InN hexagonal microdisks were grown on the LAO substrate. **Figure 5(c)** shows the tilt-view SEM image of the InN microdisk. The micrographic images of the sample showed that the 3D *c*-plane InN hexagonal microdisks and nanopillars were grown atop an anionic hexagonal basal plane of LAO, while the 2D

Compared with the shape of GaN microdisk, InN microdisk was thinner than GaN microdisk. The microstructure of the sample was analysed by field emission transmission electron microscopy (FE-TEM) (Phillips, model Tecnai F-20) with an electron voltage of 200 kV. The cross-sectional TEM specimen of the sample was prepared by a dual-beam FIB system (Seiko

thin disk. The FIB was performed with an accelerated voltage of 30 kV to cut the samples roughly and then refined the specimen further by an accelerated voltage of 5 kV. **Figure 6(a)**

2 0]

was uniquely formed by the *c*-plane wurtzite InN. The *d*-spacing between the {<sup>0</sup> <sup>0</sup> <sup>0</sup> ¯

1 0 0} planes of InN hexagonal thin disk was *d***M** = 0.3025 nm. Compared with the values on JCPDS file No. 50-1239, which are 0.5703 and 0.30647 nm, respectively, the difference between

that the lattice constant of wurtzite InN microdisk is smaller than that of bulk InN. The angle between edge and growth direction can be examined directly by the high-resolution TEM image performed at HR01 to be about 73°, as shown in **Figure 6(c)**. To establish the growth mechanism of the thin InN hexagonal microdisk, we demonstrated a ball-stick model for the self-assembled thin InN microdisk. The ball-stick model for the standard wurtzite InN (JCPDS

InN hexagonal microdisk, as shown in **Figure 6(d)** and **(e)**, where blue balls represented In atoms and red balls represented N atoms. In the case of InN thin disk, when the growth tem-

capture of N atoms by the ¯¯*β*-dangling bonds of the most-outside In atoms and then the lateral over-growth occurred; and the capture of In atoms by ¯¯*β*-dangling bonds of N atoms to form

that InN was well-formed on the InGaN buffer layer, and the InGaN buffer layer was wellestablished on the GaN epi-layer. The thicknesses of *M*-plane InN, *M*-plane InGaN and *M*plane GaN were measured to be about 265, 51 and 137 nm, respectively. The height for the *c*-plane InN hexagonal thin disk from neck to top was about 188 nm. The *c*-plane wurtzite structure was followed up to the neck area and formed a uniform *c*-plane InGaN pyramidshaped structure. Outside the pyramid-shaped structure, the wave-shaped structures were produced by the staking faults between the misfit *c*-plane wurtzite structures of InGaN and InN. The wave-shaped structures became uniform and then the *c*-plane wurtzite structure was followed further to form the InN hexagonal microdisk. The selective area diffraction (SAD) pattern at the top area of hexagonal thin disk shown in **Figure 6(b)** clearly showed one single rectangular diffraction pattern at the location of DP01, indicating that the hexagonal thin disk

2 0]

1 0 0] direction of the *c*-plane InN hexagonal

InN // [0 0 1]LAO zone axis. It clearly exhibited

= 0.5687 nm and the *d*-spacing between the

*α*/*c* = 3/8 was used to simulate the *c*-plane

1) hexagonal thin disk was built up with the

and *d***M** are 0.28 and 1.24%, respectively, revealing

InN // [0 0 1]LAO.

1} planes

*<sup>M</sup>*-plane InN epi-film was developed along with the lateral orientation [<sup>1</sup> <sup>1</sup> ¯

Inc., SII-3050), on the cleavage plane along the [1¯

of InN hexagonal thin disk was measured to be *d***<sup>c</sup>**

file No. 50-1239) with *a* = *b* = 0.3537, *c* = 0.5703 and *<sup>u</sup>* <sup>=</sup> ¯¯

perature lowered to 470°C, the *c*-plane InN (<sup>0</sup> <sup>0</sup> <sup>0</sup> ¯

wurtzite InN microdisk and bulk InN for *d***<sup>c</sup>**

{1¯

162 Epitaxy

shows the bright field image with the [<sup>1</sup> <sup>1</sup>¯

The growth mechanism of the awl-shaped GaN microdisk is divergently self-assembled, indicating that the hexagonal neck area for initial nucleation between GaN microdisk and LAO substrate is very small (diameter ~100 nm), and the strain due to the lattice-mismatch between GaN and LAO substrate will not be delivered to the awl-shaped GaN microdisk at the top. This is the way that the GaN microdisk can be grown in balance with a good awl-shape of


**Table 1.** Comparison of properties of GaN and InN microdisks.

hexagonal disk. The experimental results revealed that the awl-shaped GaN microdisk exhibited a high-quality single crystal. Therefore, the awl-shaped GaN microdisk can be regarded as a nearly freestanding substrate (strain-free) to grow the In*<sup>x</sup>* Ga1−*<sup>x</sup>* N/GaN multiple quantum wells (MQWs) on its top. The In*<sup>x</sup>* Ga1−*<sup>x</sup>* N/GaN double quantum well (DQWs) microdisk sample was grown on a high-quality 1 × 1 cm<sup>2</sup> LAO (1 0 0) substrate by low-temperature PA-MBE system (Veeco Applied-GEN 930). The LAO substrate was cut from the crystal ingot, fabricated by the traditional Czochralski pulling technique. The growth details were described completely in our previous paper which was published in *Applied Physics Letters* [18].

#### **4.2. Characteristics of InGaN/GaN microdisks**

The surface morphology of the InGaN/GaN microdisk sample was evaluated by the field emission scanning electron microscopy (FE-SEM, SII-3050). **Figure 7** shows SEM images with a tilted-angle view and a top view of the sample, respectively. The surface morphology of the sample was formed by the two-orientation growth mechanism. Comparing with the surface morphology of GaN microdisks, the shape of the as-grown InGaN/GaN DQW microdisks still maintains the hexagonal shape. **Figure 7(f)** shows the enlarged SEM image with a top view of the GaN hexagonal microdisk, which is shown in the centre of **Figure 7(d)**, and the diameter of the centre GaN microdisk is about 1.96 μm.

the FWHM value equal to (0.043 ± 0.001) eV. It is due to the band-gap transition of wurtzite

**Figure 7.** (a)–(c) The tilt-view SEM images of the sample, the scale bars are 10, 4 and 1 μm, respectively. (d)–(f) The top-

wells is higher than that from GaN due to the effect of quantum confinement. It also domi-

odoluminescence (CL, JEOL JSM-6330) and secondary electron images (SEI) measurements. The result of CL measurement is shown in **Figure 9**. We obtained the CL spectrum by detecting the photon energy from 700 (e.g. 1.77 eV) to 340 nm (e.g. 3.64 eV) with an accelerative voltage of 10 kV and an extraction voltage of photon-electric magnitude tube of 1100 V. Under such a condition, we observed two major peaks at (2.221 ± 0.0001) eV with the FWHM value equal to (0.182 ± 0.001) and (3.398 ± 0.0001) eV with the FWHM value equal to (0.046 ± 0.001) eV. We also observed the smallest peak at (2.805 ± 0.002) eV with the FWHM value equal to (0.110 ± 0.001) eV. The insets of **Figure 9** show the CL images corresponding to the three peak energies. The CL images show that the peak for the wavelength of 364 nm (e.g. 3.407 eV) is mainly emitted from wurtzite GaN microdisks. The peak for the wavelength of 560 nm (e.g.

N/GaN microdisk was used to evaluate the details of light-emitting area by cath-

N/GaN microdisks. We also observed the edge of microdisk is brighter than the central

area. It might arise from the enhanced emission by total internal reflection in the DQWs structure [20]. The optical properties of these two major peaks are consistent with the results of PL spectra, but the smallest peak for the wavelength of 450 nm (e.g. 2.756 eV) is not observed in the PL spectra. The peak at 2.805 eV was attributed to the energy level related to the structural

Ga1−*<sup>x</sup>*

Ga1−*<sup>x</sup>*

N QDWs because the defect

GaN and InN Hexagonal Microdisks http://dx.doi.org/10.5772/intechopen.70120 165

Ga1−*<sup>x</sup>*

N quantum wells in the

N quantum

GaN. The FWHM value of GaN is smaller than that of In*<sup>x</sup>*

view SEM images of the sample, the scale bars are 10, 4 and 1 μm, respectively.

nates the PL spectrum of background area tremendously.

2.214 eV) is the highest peak and mainly emitted from the In*<sup>x</sup>*

defects (e.g. Y10 in reference [15]) in microdisks.

An In*<sup>x</sup>*

In*x* Ga1−*<sup>x</sup>*

Ga1−*<sup>x</sup>*

density in GaN is much less than that in InGaN. The PL intensity from In*<sup>x</sup>*

The optical properties of the sample were measured by photoluminescence (PL, HORIBA HR800) at room temperature with a light source of He-Cd 325 nm laser. We performed the laser beam focusing on three different spots (S1–S3) and compared the results with the spot without any microdisk (i.e. mostly *M*-plane GaN and labelled as background), as shown in **Figure 8**. Two major peaks were obtained for each measurement (S1–S3). These two major peaks were confirmed by a non-linear Gaussian-function curve fitting with the software Origin (Pro. 8.0). The result of the non-linear Gaussian-function curve fitting showed that the positions of two major peaks for three spots were very consistent. The averaged value for the first peak is (2.199 ± 0.001) eV with the FWHM value equal to (0.410 ± 0.005) eV. It is due to the band-gap transition of InGaN wells. According to Vegard's law [19] with the bowing effect of bulk In*<sup>x</sup>* Ga1−*<sup>x</sup>* N: *Eg* (*x*) = [3.42 − *x*\*2.65 − *x*\*(1 − *x*)\*2.4] eV [8], we estimated the content of indium in the In*<sup>x</sup>* Ga1−*<sup>x</sup>* N DQWs, which is found to be about 28%. We note that the bowing factor needs to be modified slightly for a quantum well, and the energy-shift due to the quantum confinement will also result in a small deviation to the indium concentration of the In*<sup>x</sup>* Ga1−*<sup>x</sup>* N/GaN DQWs. The averaged value for the second peak is (3.380 ± 0.001) eV with

hexagonal disk. The experimental results revealed that the awl-shaped GaN microdisk exhibited a high-quality single crystal. Therefore, the awl-shaped GaN microdisk can be regarded

**(lateral over-growth)**

GaN microdisk 620°C One *d***M**/*d***<sup>c</sup>** 28° Base for InGaN/GaN

InN microdisk 470°C Six *d***M**/*d***<sup>c</sup>** 73° Base for InGaN/GaN

system (Veeco Applied-GEN 930). The LAO substrate was cut from the crystal ingot, fabricated by the traditional Czochralski pulling technique. The growth details were described

The surface morphology of the InGaN/GaN microdisk sample was evaluated by the field emission scanning electron microscopy (FE-SEM, SII-3050). **Figure 7** shows SEM images with a tilted-angle view and a top view of the sample, respectively. The surface morphology of the sample was formed by the two-orientation growth mechanism. Comparing with the surface morphology of GaN microdisks, the shape of the as-grown InGaN/GaN DQW microdisks still maintains the hexagonal shape. **Figure 7(f)** shows the enlarged SEM image with a top view of the GaN hexagonal microdisk, which is shown in the centre of **Figure 7(d)**, and the diameter

The optical properties of the sample were measured by photoluminescence (PL, HORIBA HR800) at room temperature with a light source of He-Cd 325 nm laser. We performed the laser beam focusing on three different spots (S1–S3) and compared the results with the spot without any microdisk (i.e. mostly *M*-plane GaN and labelled as background), as shown in **Figure 8**. Two major peaks were obtained for each measurement (S1–S3). These two major peaks were confirmed by a non-linear Gaussian-function curve fitting with the software Origin (Pro. 8.0). The result of the non-linear Gaussian-function curve fitting showed that the positions of two major peaks for three spots were very consistent. The averaged value for the first peak is (2.199 ± 0.001) eV with the FWHM value equal to (0.410 ± 0.005) eV. It is due to the band-gap transition of InGaN wells. According to Vegard's law [19] with the bowing

ing factor needs to be modified slightly for a quantum well, and the energy-shift due to the quantum confinement will also result in a small deviation to the indium concentration of

N/GaN DQWs. The averaged value for the second peak is (3.380 ± 0.001) eV with

(*x*) = [3.42 − *x*\*2.65 − *x*\*(1 − *x*)\*2.4] eV [8], we estimated the content

N DQWs, which is found to be about 28%. We note that the bow-

completely in our previous paper which was published in *Applied Physics Letters* [18].

Ga1−*<sup>x</sup>*

N/GaN double quantum well (DQWs) microdisk sam-

LAO (1 0 0) substrate by low-temperature PA-MBE

**Oblique angle Application**

QW

QW

N/GaN multiple quantum

as a nearly freestanding substrate (strain-free) to grow the In*<sup>x</sup>*

Ga1−*<sup>x</sup>*

**Growth temperature Growth mechanism** 

wells (MQWs) on its top. The In*<sup>x</sup>*

164 Epitaxy

ple was grown on a high-quality 1 × 1 cm<sup>2</sup>

**Table 1.** Comparison of properties of GaN and InN microdisks.

**4.2. Characteristics of InGaN/GaN microdisks**

of the centre GaN microdisk is about 1.96 μm.

effect of bulk In*<sup>x</sup>*

the In*<sup>x</sup>*

of indium in the In*<sup>x</sup>*

Ga1−*<sup>x</sup>*

Ga1−*<sup>x</sup>*

N: *Eg*

Ga1−*<sup>x</sup>*

**Figure 7.** (a)–(c) The tilt-view SEM images of the sample, the scale bars are 10, 4 and 1 μm, respectively. (d)–(f) The topview SEM images of the sample, the scale bars are 10, 4 and 1 μm, respectively.

the FWHM value equal to (0.043 ± 0.001) eV. It is due to the band-gap transition of wurtzite GaN. The FWHM value of GaN is smaller than that of In*<sup>x</sup>* Ga1−*<sup>x</sup>* N QDWs because the defect density in GaN is much less than that in InGaN. The PL intensity from In*<sup>x</sup>* Ga1−*<sup>x</sup>* N quantum wells is higher than that from GaN due to the effect of quantum confinement. It also dominates the PL spectrum of background area tremendously.

An In*<sup>x</sup>* Ga1−*<sup>x</sup>* N/GaN microdisk was used to evaluate the details of light-emitting area by cathodoluminescence (CL, JEOL JSM-6330) and secondary electron images (SEI) measurements. The result of CL measurement is shown in **Figure 9**. We obtained the CL spectrum by detecting the photon energy from 700 (e.g. 1.77 eV) to 340 nm (e.g. 3.64 eV) with an accelerative voltage of 10 kV and an extraction voltage of photon-electric magnitude tube of 1100 V. Under such a condition, we observed two major peaks at (2.221 ± 0.0001) eV with the FWHM value equal to (0.182 ± 0.001) and (3.398 ± 0.0001) eV with the FWHM value equal to (0.046 ± 0.001) eV. We also observed the smallest peak at (2.805 ± 0.002) eV with the FWHM value equal to (0.110 ± 0.001) eV. The insets of **Figure 9** show the CL images corresponding to the three peak energies. The CL images show that the peak for the wavelength of 364 nm (e.g. 3.407 eV) is mainly emitted from wurtzite GaN microdisks. The peak for the wavelength of 560 nm (e.g. 2.214 eV) is the highest peak and mainly emitted from the In*<sup>x</sup>* Ga1−*<sup>x</sup>* N quantum wells in the In*x* Ga1−*<sup>x</sup>* N/GaN microdisks. We also observed the edge of microdisk is brighter than the central area. It might arise from the enhanced emission by total internal reflection in the DQWs structure [20]. The optical properties of these two major peaks are consistent with the results of PL spectra, but the smallest peak for the wavelength of 450 nm (e.g. 2.756 eV) is not observed in the PL spectra. The peak at 2.805 eV was attributed to the energy level related to the structural defects (e.g. Y10 in reference [15]) in microdisks.

The microstructure of the In*<sup>x</sup>*

riers and photon emission in the In*<sup>x</sup>*

ite GaN. The *d*-spacing between the {<sup>0</sup> <sup>0</sup> <sup>0</sup>¯

Ga1−-*<sup>x</sup>*

beam FIB system (Seiko Inc., SII-3050), on the cleavage plane along the [1¯

age of 5 kV. **Figure 10(a)** and **(b)** shows the bright field image with the [<sup>1</sup> <sup>1</sup>¯

the scanning transmission electron microscopy (STEM) images of the In*<sup>x</sup>*

GaN barriers, as shown in **Figure 10(f)**, the thicknesses of In*<sup>x</sup>*

sured to be *d***<sup>c</sup>** = 0.5172 nm and the *d*-spacing between the {1¯

grown on the GaN microdisk. From the high contract image between In*<sup>x</sup>*

Ga1−*<sup>x</sup>*

transmission electron microscopy (FE-TEM) (Phillips, model Tecnai F-20) with an electron voltage of 200 kV. The cross-sectional TEM specimen of the sample was prepared by a dual-

the *c*-plane GaN hexagonal microdisk. The FIB was performed with an accelerated voltage of 30 kV to cut the samples roughly and then refined the specimen further by an accelerated volt-

zone axis, and we found that the height of the microdisk is about 2.12 μm. **Figure 10(c)** shows

evaluated from the STEM image were found to be about 2.75 and 17 nm, respectively. The quantum-well-thickness of 2.75 nm offers a good quantum confinement for the charged car-

and CL spectra. The selective area diffraction (SAD) pattern at the top area of GaN microdisk shown in **Figure 10(d)** clearly showed one single rectangular diffraction pattern at the location of DP01, indicating that the hexagonal microdisk was uniquely formed by the *c*-plane wurtz-

**Figure 10.** TEM analyses of the InGaN/GaN QW microdisk: (a) and (b) show the bright field images with [<sup>1</sup> <sup>1</sup> ¯

scale bars are 2 nm. (f) The enlarged STEM image of InGaN/GaN QW (the scale bar is 20 nm).

[0 0 1]LAO zone axis, the scale bars are 0.5 and 100 nm, respectively. (c) The STEM image taken at the top area of InGaN/ GaN QW, the scale bar is 200 nm. The selective area diffraction patterns taken at the points shown in (b) are presented in (d), the scale bar is 2 (1/nm). The high-resolution TEM images taken at the point shown in (b) are presented in (e), the

N/GaN microdisk sample was analysed by field emission

Ga1−*<sup>x</sup>*

N/GaN DQWs. It is consistent with the results of PL

1} planes of GaN hexagonal microdisk was mea-

1 0 0] direction of

167

GaN // [0 0 1]LAO

N/GaN MQWs

N wells and

2 0]GaN //

2 0]

GaN and InN Hexagonal Microdisks http://dx.doi.org/10.5772/intechopen.70120

Ga1−*<sup>x</sup>*

1 0 0} planes of GaN hexagonal

N well and GaN barrier

Ga1−*<sup>x</sup>*

**Figure 8.** The PL spectra measured at room temperature for different spots (S1–S3) of the In*<sup>x</sup>* Ga1−*<sup>x</sup>* N/GaN DQWs and the *M*-plane background of the sample.

**Figure 9.** The CL spectrum measured at room temperature for the In*<sup>x</sup>* Ga1−*<sup>x</sup>* N/GaN QW of the sample. The inset (a) shows the SEM image. The inserts (b), (c) and (d) show the CL images for wavelength of 364 nm, 433 nm, and 550 nm to the peak energies of CL spectrum.

The microstructure of the In*<sup>x</sup>* Ga1−-*<sup>x</sup>* N/GaN microdisk sample was analysed by field emission transmission electron microscopy (FE-TEM) (Phillips, model Tecnai F-20) with an electron voltage of 200 kV. The cross-sectional TEM specimen of the sample was prepared by a dualbeam FIB system (Seiko Inc., SII-3050), on the cleavage plane along the [1¯ 1 0 0] direction of the *c*-plane GaN hexagonal microdisk. The FIB was performed with an accelerated voltage of 30 kV to cut the samples roughly and then refined the specimen further by an accelerated voltage of 5 kV. **Figure 10(a)** and **(b)** shows the bright field image with the [<sup>1</sup> <sup>1</sup>¯ 2 0] GaN // [0 0 1]LAO zone axis, and we found that the height of the microdisk is about 2.12 μm. **Figure 10(c)** shows the scanning transmission electron microscopy (STEM) images of the In*<sup>x</sup>* Ga1−*<sup>x</sup>* N/GaN MQWs grown on the GaN microdisk. From the high contract image between In*<sup>x</sup>* Ga1−*<sup>x</sup>* N wells and GaN barriers, as shown in **Figure 10(f)**, the thicknesses of In*<sup>x</sup>* Ga1−*<sup>x</sup>* N well and GaN barrier evaluated from the STEM image were found to be about 2.75 and 17 nm, respectively. The quantum-well-thickness of 2.75 nm offers a good quantum confinement for the charged carriers and photon emission in the In*<sup>x</sup>* Ga1−*<sup>x</sup>* N/GaN DQWs. It is consistent with the results of PL and CL spectra. The selective area diffraction (SAD) pattern at the top area of GaN microdisk shown in **Figure 10(d)** clearly showed one single rectangular diffraction pattern at the location of DP01, indicating that the hexagonal microdisk was uniquely formed by the *c*-plane wurtzite GaN. The *d*-spacing between the {<sup>0</sup> <sup>0</sup> <sup>0</sup>¯ 1} planes of GaN hexagonal microdisk was measured to be *d***<sup>c</sup>** = 0.5172 nm and the *d*-spacing between the {1¯ 1 0 0} planes of GaN hexagonal

**Figure 8.** The PL spectra measured at room temperature for different spots (S1–S3) of the In*<sup>x</sup>*

**Figure 9.** The CL spectrum measured at room temperature for the In*<sup>x</sup>*

peak energies of CL spectrum.

Ga1−*<sup>x</sup>*

the SEM image. The inserts (b), (c) and (d) show the CL images for wavelength of 364 nm, 433 nm, and 550 nm to the

*M*-plane background of the sample.

166 Epitaxy

Ga1−*<sup>x</sup>*

N/GaN QW of the sample. The inset (a) shows

N/GaN DQWs and the

**Figure 10.** TEM analyses of the InGaN/GaN QW microdisk: (a) and (b) show the bright field images with [<sup>1</sup> <sup>1</sup> ¯ 2 0]GaN // [0 0 1]LAO zone axis, the scale bars are 0.5 and 100 nm, respectively. (c) The STEM image taken at the top area of InGaN/ GaN QW, the scale bar is 200 nm. The selective area diffraction patterns taken at the points shown in (b) are presented in (d), the scale bar is 2 (1/nm). The high-resolution TEM images taken at the point shown in (b) are presented in (e), the scale bars are 2 nm. (f) The enlarged STEM image of InGaN/GaN QW (the scale bar is 20 nm).

microdisk was *d***M** = 0.2764 nm. The angle between edge and growth direction can be obtained by *φ* = tan−1(*d***M**/*d***<sup>c</sup>** ) = 28.1°, which was in good agreement with the model predicted. The angle between the edge and the growth direction examined directly by the STEM image can be found about 28°, as shown in **Figure 10(c)**. The high-quality crystalline micro structures for the In*<sup>x</sup>* Ga1−*<sup>x</sup>* N/GaN DQWs were reconfirmed by the high-resolution TEM images, as shown in **Figure 10(e)**. It showed that the In*<sup>x</sup>* Ga1−*<sup>x</sup>* N wells were well-stacked on the high-quality GaN microdisk to form well-assembled crystalline In*<sup>x</sup>* Ga1−*<sup>x</sup>* N/GaN DQWs with some minor structural defects (e.g. dislocations or stacking-faults) occurred in the In*<sup>x</sup>* Ga1−*<sup>x</sup>* N wells. The HR results indicated that the density of structural defects in In*<sup>x</sup>* Ga1−*<sup>x</sup>* N wells is greater than that in GaN layers. It supported the results of FWHM analyses of PL spectra, indicating that FWHM (0.173 eV) for the peak of 2.192 eV from In*<sup>x</sup>* Ga1−*<sup>x</sup>* N wells was greater than FWHM (0.043 eV) for the peak of 3.383 eV from the wurtzite GaN layer. The structural defeats in In*<sup>x</sup>* Ga1−*<sup>x</sup>* N wells yield the peak of 2.805 eV in CL measurement.

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## **5. Conclusion**

We have grown GaN and InN hexagonal microdisks on the LAO substrates at low temperatures (GaN at 630°C and InN at 470°C) by PA-MBE. From the SEM images and TEM analyses, we found that 3D *c*-plane hexagonal microdisks and 2D *M*-plane epi-film were grown on the LAO substrate. From TEM analyses, the oblique angle of GaN and InN hexagonal microdisks can be examined directly to be about 28 and 73°, respectively. The lateral over-growth mechanism causes the bigger oblique angle of InN hexagonal thin disks at low growth temperature. From PL and CL analyses, we found that the high-intensity light of 367-nm wavelength (3.380 eV) and 566-nm wavelength (2.192 eV) emitted from the GaN microdisks and In*<sup>x</sup>* Ga1−*<sup>x</sup>* N/GaN quantum wells, respectively. Therefore, the stain-free microdisk provides an opportunity to fabricate In*x* Ga1−*<sup>x</sup>* N/GaN microdisk quantum well for the application of full-colour micron LED without the sapphire substrate.

## **Acknowledgements**

The project was supported by the Ministry of Science and Technology of Taiwan and the Core Facilities Laboratory for Nanoscience and Nanotechnology in Kaohsiung and Pintung Area.

## **Author details**

Chen-Chi Yang, Ikai Lo\*, Yu-Chi Hsu and Hong-Yi Yang

\*Address all correspondence to: ikailo@mail.phys.nsysu.edu.tw

Department of Physics, Center for Nanoscience and Nanotechnology, National Sun Yat-Sen University, Kaohsiung, Taiwan

## **References**

microdisk was *d***M** = 0.2764 nm. The angle between edge and growth direction can be obtained

between the edge and the growth direction examined directly by the STEM image can be found about 28°, as shown in **Figure 10(c)**. The high-quality crystalline micro structures for

GaN layers. It supported the results of FWHM analyses of PL spectra, indicating that FWHM

Ga1−*<sup>x</sup>*

We have grown GaN and InN hexagonal microdisks on the LAO substrates at low temperatures (GaN at 630°C and InN at 470°C) by PA-MBE. From the SEM images and TEM analyses, we found that 3D *c*-plane hexagonal microdisks and 2D *M*-plane epi-film were grown on the LAO substrate. From TEM analyses, the oblique angle of GaN and InN hexagonal microdisks can be examined directly to be about 28 and 73°, respectively. The lateral over-growth mechanism causes the bigger oblique angle of InN hexagonal thin disks at low growth temperature. From PL and CL analyses, we found that the high-intensity light of 367-nm wavelength (3.380 eV)

tum wells, respectively. Therefore, the stain-free microdisk provides an opportunity to fabricate

The project was supported by the Ministry of Science and Technology of Taiwan and the Core Facilities Laboratory for Nanoscience and Nanotechnology in Kaohsiung and Pintung Area.

Department of Physics, Center for Nanoscience and Nanotechnology, National Sun Yat-Sen

N/GaN microdisk quantum well for the application of full-colour micron LED without

Ga1−*<sup>x</sup>*

for the peak of 3.383 eV from the wurtzite GaN layer. The structural defeats in In*<sup>x</sup>*

and 566-nm wavelength (2.192 eV) emitted from the GaN microdisks and In*<sup>x</sup>*

Chen-Chi Yang, Ikai Lo\*, Yu-Chi Hsu and Hong-Yi Yang

\*Address all correspondence to: ikailo@mail.phys.nsysu.edu.tw

tural defects (e.g. dislocations or stacking-faults) occurred in the In*<sup>x</sup>*

) = 28.1°, which was in good agreement with the model predicted. The angle

N wells were well-stacked on the high-quality GaN

Ga1−*<sup>x</sup>*

N/GaN DQWs with some minor struc-

Ga1−*<sup>x</sup>*

N wells was greater than FWHM (0.043 eV)

N wells is greater than that in

Ga1−*<sup>x</sup>*

N/GaN quan-

N wells. The HR

Ga1−*<sup>x</sup>*

N wells

N/GaN DQWs were reconfirmed by the high-resolution TEM images, as shown in

Ga1−*<sup>x</sup>*

by *φ* = tan−1(*d***M**/*d***<sup>c</sup>**

Ga1−*<sup>x</sup>*

**5. Conclusion**

the sapphire substrate.

**Acknowledgements**

**Author details**

University, Kaohsiung, Taiwan

In*x* Ga1−*<sup>x</sup>*

**Figure 10(e)**. It showed that the In*<sup>x</sup>*

microdisk to form well-assembled crystalline In*<sup>x</sup>*

(0.173 eV) for the peak of 2.192 eV from In*<sup>x</sup>*

yield the peak of 2.805 eV in CL measurement.

results indicated that the density of structural defects in In*<sup>x</sup>*

the In*<sup>x</sup>*

168 Epitaxy


[14] Lo I, Hsieh CH, Chen YL, Pang WY, Hsu YC,Chiang JC, Chou MC, Tsai JK, Schaadt DM. Line defects of M-plane GaN grown on g-LiAlO<sup>2</sup> by plasma-assisted molecular beam epitaxy. Applied Physics Letters. 2008;**92**:202106. DOI: 10.1063/1.2924288

**Chapter 8**

Provisional chapter

**Heterostructures of III-Nitride Semiconductors for**

DOI: 10.5772/intechopen.70219

Heterostructures of III-Nitride Semiconductors for

III-Nitride-based heterostructures are well suited for the fabrication of various optoelectronic devices such as light-emitting diodes (LEDs), laser diodes (LDs), high-power/ high-frequency field-effect transistors (FETs), and tandem solar cells because of their inherent properties. However, the heterostructures grown along polar direction are affected by the presence of internal electric field induced by the existence of intrinsic spontaneous and piezoelectric polarizations. The internal electric field is deleterious for optoelectronic devices as it causes a spatial separation of electron and hole wave functions in the quantum wells, which thereby decreases the emission efficiency. The growth of III-nitride heterostructures in nonpolar or semipolar directions is an alternative option to minimize the piezoelectric polarization. The heterostructures grown on these orientations are receiving a lot of focus due to their potential improvement on the efficiency of optoelectronic devices. In the present chapter, the growth of polar and nonpolar III-nitride heterostructures using molecular beam epitaxy (MBE) system and their characterizations are discussed. The transport properties of the III-nitride heterostructure-based Schottky junctions are also included. In addition, their applications toward UV and IR

Keywords: heterostructures, InN/GaN, InGaN/Si, InGaN/GaN, HRXRD

There has been remarkable progress in the development of group III-nitride-based heterostructures because of their potential application in fabricating various optical and electronic devices such as light-emitting diodes (LEDs), laser diodes (LDs), tandem solar cells, field-effect transistors (FETs), and Schottky junctions. Heterostructures are ubiquitous of semiconductor

> © The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and eproduction in any medium, provided the original work is properly cited.

© 2018 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

**Optical and Electronic Applications**

Optical and Electronic Applications

Shruti Mukundan and Saluru Baba Krupanidhi

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

Basanta Roul, Greeshma Chandan, Shruti Mukundan and Saluru Baba Krupanidhi

Basanta Roul, Greeshma Chandan,

http://dx.doi.org/10.5772/intechopen.70219

Abstract

detectors are discussed.

1. Introduction


Provisional chapter

## **Heterostructures of III-Nitride Semiconductors for Optical and Electronic Applications** Heterostructures of III-Nitride Semiconductors for

DOI: 10.5772/intechopen.70219

Basanta Roul, Greeshma Chandan, Shruti Mukundan and Saluru Baba Krupanidhi Basanta Roul, Greeshma Chandan, Shruti

Optical and Electronic Applications

Additional information is available at the end of the chapter Mukundan and Saluru Baba Krupanidhi

http://dx.doi.org/10.5772/intechopen.70219 Additional information is available at the end of the chapter

#### Abstract

[14] Lo I, Hsieh CH, Chen YL, Pang WY, Hsu YC,Chiang JC, Chou MC, Tsai JK, Schaadt DM.

[15] Reshchikov MA, Morkoc HJ. Luminescence properties of defects in GaN. Journal of

[16] Yang CC, Lo I, Hu CH, Huang HC, Chou MMC. Growth of InN hexagonal microdisks.

[17] Denton AR, Ashcroft NW. Vegar's law. Physical Review A. 1991;**43**(6):3161. DOI: 10.1103/

[18] Hsu YC, Lo I, Shih CH, Pang WY, Hu CH, Wang YC, Tsai CD, Chou MMC, Gary Z. Green light emission by InGaN/GaN multiple-quantum-well microdisks. Applied

[19] Vegard L. Die Konstitution der Mischkristalle und die Raumfüllung der Atome.

[20] Choi S, Ton-That C, Phillips MR, Aharonovich I. Observation of whispering gallery modes from hexagonal ZnO microdisks using cathodoluminescence spectroscopy.

epitaxy. Applied Physics Letters. 2008;**92**:202106. DOI: 10.1063/1.2924288

by plasma-assisted molecular beam

Line defects of M-plane GaN grown on g-LiAlO<sup>2</sup>

Applied Physics. 2005;**97**:061301. DOI: 10.1063/1.1868059

AIP Advances. 2015;**6**:085015. DOI: 10.1063/1.4961699

Physics Letters. 2014;**104**:102105. DOI: 10.1063/1.4868417

Zeitschrift Fur Physik. 1921;**5**:17-26. DOI: 10.1007/BF01349680

Applied Physics Letters. 2013;**103**:171102. DOI: 10.1063/1.482648s1

PhysRevA.43.3161

170 Epitaxy

III-Nitride-based heterostructures are well suited for the fabrication of various optoelectronic devices such as light-emitting diodes (LEDs), laser diodes (LDs), high-power/ high-frequency field-effect transistors (FETs), and tandem solar cells because of their inherent properties. However, the heterostructures grown along polar direction are affected by the presence of internal electric field induced by the existence of intrinsic spontaneous and piezoelectric polarizations. The internal electric field is deleterious for optoelectronic devices as it causes a spatial separation of electron and hole wave functions in the quantum wells, which thereby decreases the emission efficiency. The growth of III-nitride heterostructures in nonpolar or semipolar directions is an alternative option to minimize the piezoelectric polarization. The heterostructures grown on these orientations are receiving a lot of focus due to their potential improvement on the efficiency of optoelectronic devices. In the present chapter, the growth of polar and nonpolar III-nitride heterostructures using molecular beam epitaxy (MBE) system and their characterizations are discussed. The transport properties of the III-nitride heterostructure-based Schottky junctions are also included. In addition, their applications toward UV and IR detectors are discussed.

Keywords: heterostructures, InN/GaN, InGaN/Si, InGaN/GaN, HRXRD

## 1. Introduction

There has been remarkable progress in the development of group III-nitride-based heterostructures because of their potential application in fabricating various optical and electronic devices such as light-emitting diodes (LEDs), laser diodes (LDs), tandem solar cells, field-effect transistors (FETs), and Schottky junctions. Heterostructures are ubiquitous of semiconductor

© 2018 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

© The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and eproduction in any medium, provided the original work is properly cited.

devices, and most of semiconductor devices have two or more semiconductor materials. A heterostructure is formed between two layers of dissimilar semiconductors having unequal energy bandgap. In order to realize high-performance devices, growth of device quality heterostructures is required. Yoshida et al. reported improved cathodoluminescence efficiency of the GaN layer grown on sapphire using AlN as a buffer layer [1]. Later Akasaki et al. [2] and Nakamura et al. [3] had employed a two-step growth method, where a nucleation layer was grown at low temperature followed by the GaN layer at high temperatures. Nakamura et al. fabricated the first blue LED consisting of a p-GaN/n-InGaN/n-GaN double heterostructure in 1993 [4] for which he won the Nobel Prize in 2014 and the first violet laser consisting of InGaN/ GaN/AlGaN heterostructure in 1996 [5]. Similarly, Khan et al. [6] achieved the first breakthrough in the field of high mobility transistors based on AlGaN/GaN heterostructure in 1994.

So far most of indium gallium nitride (InGaN)-based LEDs are built along Ga-polar (0001) orientation, which is susceptible to the strong internal electric field induced by the spontaneous and the piezoelectric polarization in wurtzite III-nitrides. The effect of polarization will be explained in detail in the coming section. There have been concerted efforts in exploring IIInitride materials and devices along nonpolar and semipolar crystallographic orientations [7, 8]. The two major challenges in the field of InGaN-based LEDs are the "efficiency droop" under a high injection current density and the "green gap" in the plot of efficiency versus emission wavelength [9]. Very promising reports of LEDs and laser diodes on nonpolar and semipolar GaN bulk substrates, in the longer wavelength of green and yellow, tend to validate the concept of crystallographic engineering [10]. However, nonpolar and semipolar GaN bulk substrates are presently very small in size and expensive in price. Therefore, most of the current research in this field is focused on the growth of high-quality epilayers on nonnative substrates which are available in large wafer sizes and also cost-effective than the native substrates, thus paving the way for the commercialization of devices based on nonpolar and semipolar GaN [11]. However, the lattice mismatch along different directions poses difficulty in the hetero-epitaxial growth of nonpolar (a- and m-plane) and semipolar GaN on the foreign substrates, often resulting in nonuniform nucleation. This leads to the growth of GaN with a defective microstructure, arising due to the formation of basal-plane stacking faults (BSFs) and partial dislocations (PDs) [12, 13].

presence of external stress due to lattice mismatch in the films grown on foreign substrates, or in heterostructures, results in additional, piezoelectric contribution to the polarization. The

Figure 1. Ball and stick model illustrating crystal structure of wurtzite Ga-polarity and N-polarity GaN. Reprinted with

The type of stress determines the direction of piezoelectric polarization. When a layer of smaller lattice constant than that of GaN, such as AlGaN, is grown on GaN, the grown AlGaN layer experiences tensile stress. Whereas when a layer with larger lattice constant than GaN such as InGaN is grown on GaN, the resulting strain in the InGaN layer is compressive. The piezoelectric field generated as a result of the tensile strain in the case of AlGaN/GaN and compressive strain in the case of InGaN/GaN is parallel and antiparallel, respectively. The overall polarization effect in InGaN is therefore smaller compared to AlGaN [16]. The polarization field plays a very pivotal role in GaN-based LED devices. Currently, all highly efficient blue or UV LED devices are based on multi-quantum well (MQW) structures. In a MQW structure, very few atomic layers of a narrow bandgap material, referred to as wells, are sandwiched between thicker wide bandgap materials, referred to as barriers. In quantum well (QW) structures, the charge carriers are confined to wells with high-energy barriers on either side, thus preventing the charge carriers from escaping without recombining with their counterparts and thus increasing the probability of radiative recombination. The presence of spontaneous and piezoelectric polarization in QW leads to asymmetry in the electric-field profile and results in bending of the conduction and valence bands, thus spatially separating the charge carriers. Due to this spatial separation of the charge carriers, the overlapping of charge

Ptotal ¼ Pspon þ Ppiezo (1)

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total polarization, therefore, can be presented as the sum of the two components:

permission from [15].

The absence of inversion symmetry in wurtzite GaN along the [0001] direction indicates that [0001] and [0001] directions are not equivalent. Along this crystal direction if one face ends with Ga atoms, the other face will end with N atoms in the place of Ga atom and vice versa. This asymmetric arrangement of Ga and N atoms along the [0001] direction gives rise to charge polarity and thus can be referred as microscopic dipoles. The presence of polarity and lack of inversion symmetry lead to the generation of macroscopic polarization along [0001] direction and are usually referred to as spontaneous polarization. The word "spontaneous" means that it is arising only due to the crystallographic arrangement of the end faces but not due to the strain [14]. In III-nitrides, asymmetry of inversion is present only along the c-axis. Hence, PSP is parallel to this direction, and c-plane nitrides are therefore called polar nitride materials. The c-axis direction consists of two opposite stacking sequences of atomic layering, resulting in either cation-face (metal) or anion-face (nitrogen) epitaxy. Ga and In polarity is in the [0001] direction, and N-polarity is in the [0001] direction as shown in Figure 1 [15]. The

devices, and most of semiconductor devices have two or more semiconductor materials. A heterostructure is formed between two layers of dissimilar semiconductors having unequal energy bandgap. In order to realize high-performance devices, growth of device quality heterostructures is required. Yoshida et al. reported improved cathodoluminescence efficiency of the GaN layer grown on sapphire using AlN as a buffer layer [1]. Later Akasaki et al. [2] and Nakamura et al. [3] had employed a two-step growth method, where a nucleation layer was grown at low temperature followed by the GaN layer at high temperatures. Nakamura et al. fabricated the first blue LED consisting of a p-GaN/n-InGaN/n-GaN double heterostructure in 1993 [4] for which he won the Nobel Prize in 2014 and the first violet laser consisting of InGaN/ GaN/AlGaN heterostructure in 1996 [5]. Similarly, Khan et al. [6] achieved the first breakthrough

in the field of high mobility transistors based on AlGaN/GaN heterostructure in 1994.

partial dislocations (PDs) [12, 13].

172 Epitaxy

So far most of indium gallium nitride (InGaN)-based LEDs are built along Ga-polar (0001) orientation, which is susceptible to the strong internal electric field induced by the spontaneous and the piezoelectric polarization in wurtzite III-nitrides. The effect of polarization will be explained in detail in the coming section. There have been concerted efforts in exploring IIInitride materials and devices along nonpolar and semipolar crystallographic orientations [7, 8]. The two major challenges in the field of InGaN-based LEDs are the "efficiency droop" under a high injection current density and the "green gap" in the plot of efficiency versus emission wavelength [9]. Very promising reports of LEDs and laser diodes on nonpolar and semipolar GaN bulk substrates, in the longer wavelength of green and yellow, tend to validate the concept of crystallographic engineering [10]. However, nonpolar and semipolar GaN bulk substrates are presently very small in size and expensive in price. Therefore, most of the current research in this field is focused on the growth of high-quality epilayers on nonnative substrates which are available in large wafer sizes and also cost-effective than the native substrates, thus paving the way for the commercialization of devices based on nonpolar and semipolar GaN [11]. However, the lattice mismatch along different directions poses difficulty in the hetero-epitaxial growth of nonpolar (a- and m-plane) and semipolar GaN on the foreign substrates, often resulting in nonuniform nucleation. This leads to the growth of GaN with a defective microstructure, arising due to the formation of basal-plane stacking faults (BSFs) and

The absence of inversion symmetry in wurtzite GaN along the [0001] direction indicates that [0001] and [0001] directions are not equivalent. Along this crystal direction if one face ends with Ga atoms, the other face will end with N atoms in the place of Ga atom and vice versa. This asymmetric arrangement of Ga and N atoms along the [0001] direction gives rise to charge polarity and thus can be referred as microscopic dipoles. The presence of polarity and lack of inversion symmetry lead to the generation of macroscopic polarization along [0001] direction and are usually referred to as spontaneous polarization. The word "spontaneous" means that it is arising only due to the crystallographic arrangement of the end faces but not due to the strain [14]. In III-nitrides, asymmetry of inversion is present only along the c-axis. Hence, PSP is parallel to this direction, and c-plane nitrides are therefore called polar nitride materials. The c-axis direction consists of two opposite stacking sequences of atomic layering, resulting in either cation-face (metal) or anion-face (nitrogen) epitaxy. Ga and In polarity is in the [0001] direction, and N-polarity is in the [0001] direction as shown in Figure 1 [15]. The

Figure 1. Ball and stick model illustrating crystal structure of wurtzite Ga-polarity and N-polarity GaN. Reprinted with permission from [15].

presence of external stress due to lattice mismatch in the films grown on foreign substrates, or in heterostructures, results in additional, piezoelectric contribution to the polarization. The total polarization, therefore, can be presented as the sum of the two components:

$$P\_{\text{total}} = P\_{\text{spon}} + P\_{\text{piezo}} \tag{1}$$

The type of stress determines the direction of piezoelectric polarization. When a layer of smaller lattice constant than that of GaN, such as AlGaN, is grown on GaN, the grown AlGaN layer experiences tensile stress. Whereas when a layer with larger lattice constant than GaN such as InGaN is grown on GaN, the resulting strain in the InGaN layer is compressive. The piezoelectric field generated as a result of the tensile strain in the case of AlGaN/GaN and compressive strain in the case of InGaN/GaN is parallel and antiparallel, respectively. The overall polarization effect in InGaN is therefore smaller compared to AlGaN [16]. The polarization field plays a very pivotal role in GaN-based LED devices. Currently, all highly efficient blue or UV LED devices are based on multi-quantum well (MQW) structures. In a MQW structure, very few atomic layers of a narrow bandgap material, referred to as wells, are sandwiched between thicker wide bandgap materials, referred to as barriers. In quantum well (QW) structures, the charge carriers are confined to wells with high-energy barriers on either side, thus preventing the charge carriers from escaping without recombining with their counterparts and thus increasing the probability of radiative recombination. The presence of spontaneous and piezoelectric polarization in QW leads to asymmetry in the electric-field profile and results in bending of the conduction and valence bands, thus spatially separating the charge carriers. Due to this spatial separation of the charge carriers, the overlapping of charge carrier wave functions is substantially reduced resulting in lower recombination probability. The bending of bands also leads to bandgap shrinkage. As a result the emitted radiation is redshifted [17]. This process is referred to as quantum-confined Stark effect and is undesirable.

c-axis, which have equal number of Group III and V atoms and are called nonpolar surfaces. Alternatively, inclined surfaces such as (1 0 1 3), (1 0 1 1), and (1 1 2 2) are known to

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Compared to the conventionally oriented c-plane GaN, nonpolar and semipolar planes were considered "unstable" for a long time. Rough and faceted surfaces have been a prolonged problem for device fabrication on these directions [21, 22]. In 2000, Waltereit et al. first demonstrated planer m-plane GaN growth via molecular beam epitaxy [7]. This demonstration was followed by Craven et al.'s metal organic chemical vapor deposition growth of a-plane GaN films in 2002 [23]. While considerable progress started after this breakthrough in the growth of thin films in nonpolar orientation, thick-film or bulk growth of this orientation continued to be elusive, hence limiting the performance of nonpolar GaN-based devices due to lack of suitable substrates. Achieving low defect density, nonpolar oriented film is a trending problem and exponential progress is seen in this field over the past decade. In the present chapter, the growth of polar and nonpolar III-nitride heterostructures using molecular beam epitaxy (MBE) and their characterizations will be presented. The transport properties of the III-nitride heterostructure-based Schottky junctions will be presented. In addition, applications toward

have lower polarization fields and are often called semipolar planes [20].

2. Growth and characterizations of III-nitride heterostructures

The InN/GaN heterostructure system has several advantages which includes the high rate of

1

tion of high-quality InN/GaN heterostructures is a challenging issue due to the difficulty in continuous growth of InN and GaN films because of large difference in the optimum growth temperature between them (InN ~500C and GaN ~750C). In addition, due to the low dissociation temperature of InN, it is very difficult to get high-quality InN/GaN heterostructure at the growth temperature of InN. In addition, the large lattice mismatch between InN and GaN (11%) results in a poor interface. Several groups [24–26] have studied the interfaces of the heterostructures like InN/GaN, GaN/ZnO, and GaN/AlN. InN/GaN MQWs with 1 and fractional monolayers of InN were proposed and experimentally demonstrated by Yoshikawa et al. [27]. Similarly, InN/GaN single-quantum well and double heterostructures were fabricated by the PAMBE [28]. In this section, InN/GaN heterostructures were grown using MBE system. InN thin films of thicknesses around 300 nm were epitaxially grown on 4 μm-GaN/ Al2O3 (0 0 0 1) templates at different substrate temperatures. Thermal cleaning of the GaN templates was carried out at 700C for 5 min in the presence of nitrogen plasma. Following that, a step growth method was employed to grow high-quality InN epilayers. The first step involved the growth of low-temperature InN nucleation layer at 400C for 15 min, which resulted in the formation of thin buffer layer ~20 nm. Subsequently, the substrate temperature was increased to 450C (sample A), 470C (sample B), 500C (sample C), and 530C (sample D)

), high peak value of the steady-state electron

), and large conduction band offset. However, the fabrica-

UV and IR detectors will be discussed.

2.1. Polar InN/GaN heterostructures

drift velocity in InN (5 107 cm s<sup>1</sup>

optical phonon emission in InN (2.5 1013 <sup>s</sup>

Apart from the above-mentioned detrimental effects, the effective width of QWs is reduced as a consequence of tilt in the band edges which leads to a higher charge carrier density and may eventually lead to nonradiative Auger recombination. The effective barrier height is also lowered due to band bending, which means the carrier confinement is weakened with increasing bias voltage leading to carrier leakage. It is believed that these two non-radiative recombination mechanisms are responsible for the reduced efficiency in GaN-based LEDs when operated at higher current [18]. Besides wurtzite nitrides, cubic nitrides also have similar bandgaps and are free from the spontaneous polarization. However, due to the instability of cubic nitrides and the poor quality, they are less preferred for device applications. Growth of wurtzite materials that have either no polarization field or reduced one in the growth direction can solve these problems. Figure 2(b–i) shows the wurtzite III-nitride planes which are perpendicular or inclined to the [0001] direction [19]. There are two surfaces perpendicular to the

Figure 2. Schematic views of (a) polar c-plane; (b) and (c) nonpolar a- and m-planes; (d)–(i) various semipolar s-planes. Reprinted with permission from [19].

c-axis, which have equal number of Group III and V atoms and are called nonpolar surfaces. Alternatively, inclined surfaces such as (1 0 1 3), (1 0 1 1), and (1 1 2 2) are known to have lower polarization fields and are often called semipolar planes [20].

Compared to the conventionally oriented c-plane GaN, nonpolar and semipolar planes were considered "unstable" for a long time. Rough and faceted surfaces have been a prolonged problem for device fabrication on these directions [21, 22]. In 2000, Waltereit et al. first demonstrated planer m-plane GaN growth via molecular beam epitaxy [7]. This demonstration was followed by Craven et al.'s metal organic chemical vapor deposition growth of a-plane GaN films in 2002 [23]. While considerable progress started after this breakthrough in the growth of thin films in nonpolar orientation, thick-film or bulk growth of this orientation continued to be elusive, hence limiting the performance of nonpolar GaN-based devices due to lack of suitable substrates. Achieving low defect density, nonpolar oriented film is a trending problem and exponential progress is seen in this field over the past decade. In the present chapter, the growth of polar and nonpolar III-nitride heterostructures using molecular beam epitaxy (MBE) and their characterizations will be presented. The transport properties of the III-nitride heterostructure-based Schottky junctions will be presented. In addition, applications toward UV and IR detectors will be discussed.

## 2. Growth and characterizations of III-nitride heterostructures

#### 2.1. Polar InN/GaN heterostructures

carrier wave functions is substantially reduced resulting in lower recombination probability. The bending of bands also leads to bandgap shrinkage. As a result the emitted radiation is redshifted [17]. This process is referred to as quantum-confined Stark effect and is undesirable.

Apart from the above-mentioned detrimental effects, the effective width of QWs is reduced as a consequence of tilt in the band edges which leads to a higher charge carrier density and may eventually lead to nonradiative Auger recombination. The effective barrier height is also lowered due to band bending, which means the carrier confinement is weakened with increasing bias voltage leading to carrier leakage. It is believed that these two non-radiative recombination mechanisms are responsible for the reduced efficiency in GaN-based LEDs when operated at higher current [18]. Besides wurtzite nitrides, cubic nitrides also have similar bandgaps and are free from the spontaneous polarization. However, due to the instability of cubic nitrides and the poor quality, they are less preferred for device applications. Growth of wurtzite materials that have either no polarization field or reduced one in the growth direction can solve these problems. Figure 2(b–i) shows the wurtzite III-nitride planes which are perpendicular or inclined to the [0001] direction [19]. There are two surfaces perpendicular to the

Figure 2. Schematic views of (a) polar c-plane; (b) and (c) nonpolar a- and m-planes; (d)–(i) various semipolar s-planes.

Reprinted with permission from [19].

174 Epitaxy

The InN/GaN heterostructure system has several advantages which includes the high rate of optical phonon emission in InN (2.5 1013 <sup>s</sup> 1 ), high peak value of the steady-state electron drift velocity in InN (5 107 cm s<sup>1</sup> ), and large conduction band offset. However, the fabrication of high-quality InN/GaN heterostructures is a challenging issue due to the difficulty in continuous growth of InN and GaN films because of large difference in the optimum growth temperature between them (InN ~500C and GaN ~750C). In addition, due to the low dissociation temperature of InN, it is very difficult to get high-quality InN/GaN heterostructure at the growth temperature of InN. In addition, the large lattice mismatch between InN and GaN (11%) results in a poor interface. Several groups [24–26] have studied the interfaces of the heterostructures like InN/GaN, GaN/ZnO, and GaN/AlN. InN/GaN MQWs with 1 and fractional monolayers of InN were proposed and experimentally demonstrated by Yoshikawa et al. [27]. Similarly, InN/GaN single-quantum well and double heterostructures were fabricated by the PAMBE [28]. In this section, InN/GaN heterostructures were grown using MBE system. InN thin films of thicknesses around 300 nm were epitaxially grown on 4 μm-GaN/ Al2O3 (0 0 0 1) templates at different substrate temperatures. Thermal cleaning of the GaN templates was carried out at 700C for 5 min in the presence of nitrogen plasma. Following that, a step growth method was employed to grow high-quality InN epilayers. The first step involved the growth of low-temperature InN nucleation layer at 400C for 15 min, which resulted in the formation of thin buffer layer ~20 nm. Subsequently, the substrate temperature was increased to 450C (sample A), 470C (sample B), 500C (sample C), and 530C (sample D) to grow active InN epilayers. The nitrogen flow rate and the forward RF power of the plasma source were set to 0.5 sccm and 350 W. The indium BEP was 1.53 <sup>10</sup><sup>7</sup> mbar.

Figure 3(a) shows a 2θ-ω scan of InN films grown at different growth temperatures [29]. The peaks at 2θ = 31.3 and 65.5 are attributed to the (0002) and (0004) planes of the InN epilayers, whereas the peaks at 2θ = 34.56 and 72.81 are attributed to the (0002) and (0004) planes of the GaN templates. The peak at 2θ = 41.69 corresponds to the (0006) plane of sapphire substrate. The InN films grown at low (450C) and high (530C) temperatures show a peak at around 2θ = 33 indicating the presence of In metal. The presence of In metal could be due to the low migration of In at low temperature and more dissociation of InN at high temperature. The structural quality of the films was evaluated from the full width at half maximum (FWHM) of (0002) InN X-ray rocking curve (XRC). The rocking curves of the (0002) InN reflection are shown in Figure 3(b) [29]. Growth of InN epilayers at high temperature improved crystal quality, i.e., the FWHM of the (0002) InN XRC decreased from 532.8 to 450 arcsec corresponding to growth temperatures 450 and 500C. The low migration velocities of In and N atoms at low growth temperature are the most probable reason for the relatively inferior crystal quality [30]. On the other hand, growth of InN at higher temperature (530C) resulted in pronounced dissociation of InN, thus leading to high FWHM of 716.4 arcsec. The screw dislocation density of InN films, as calculated from the FWHM of the rocking curves, was found to be 2.27 108 , 2.03 108 , 1.62 108 , and 4.12 108 cm<sup>2</sup> for samples A, B, C, and D, respectively. In addition, InN films show n-type conductivity with carrier concentrations in the order of 1018 to 1019 cm<sup>3</sup> . As growth temperature increases from 450 to 500C, the carrier concentration decreases from 1.57 1019 to 3.10 1018 cm<sup>3</sup> . The high carrier concentration at low growth temperature may be due to relatively poor crystallinity. On the other hand, the carrier concentration increased to 1.10 1019 cm<sup>3</sup> at high growth temperature of 530C, which may be associated with the large dissociation of InN.

The room temperature optical absorption spectra squared as a function of growth temperature is shown in Figure 3(c) [29]. The absorption spectrum exhibits characteristic interference fringes due to the underlying thicker GaN epilayer. The absorption edge of the InN films was estimated by extrapolating the linear part of the squared absorption down to the photon energy axis. The energy corresponding to this absorption edge is the amount of energy needed by an electron to make a vertical transition from the upper valence band to the Fermi surface in the conduction band. Therefore, this energy can be considered as the Fermi-level energy in the conduction band. A strong shift was observed in the absorption edge with a change in carrier concentration. This is usually referred as Burstein-Moss shift which is a commonly observed phenomenon in narrow-gap semiconductors owing to non-parabolic conduction band [31].

there are very few reports on the growth of nonpolar InN films [36, 37]. The earlier reports on the InN on r-sapphire substrates indicate the growth of cubic (001) [38, 39] and polar (0001) InN [40]. The nonpolar a-plane InN was demonstrated by using GaN buffer layer on r-plane sapphire [41, 42]. In this section, nonpolar (1 1 2 0) a-InN/GaN heterostructures were grown on rplane and m-plane sapphire substrates, respectively, using MBE system. Prior to the growth, thermal cleaning of the sapphire substrate was carried out at 850C inside MBE chamber for 30 min under ultrahigh vacuum. RF power and flow rate were kept constant at 350 W and 0.5 sccm, respectively. (1 1 2 0) a-GaN buffer layer was grown at 760C temperature. The growth of InN epilayers was carried out using two-step growth processes: growth of low-temperature InN buffer layer (~20 nm) on GaN under layer at 400C and growth of InN epilayers at different

Figure 3. (a) 2θ-ω HRXRD scanning curve of InN films grown at different growth temperatures, (b) the XRC of the (0002) reflections of InN films grown at different growth temperatures, and (c) optical absorption spectra of InN films. Reprinted

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with permission from [29].

temperatures. The growth temperature of (1 1 2 0) a-InN varied from 470 to 530C.

Figure 4(a) show 2θ-ω scans of nonpolar (1 1 2 0) a-plane InN/GaN heterostructures grown on r-plane sapphire substrate at different growth temperatures [42]. The peaks at 2θ = 57.7 and 51.61 are assigned to the a-plane GaN and a-plane InN reflections, respectively, along with that from the r-plane sapphire substrate. The sample grown at 530C shows the peak at 32.9 which corresponds to In (101) reflections. Figure 4(b) shows the RHEED patterns in the azimuths [0001] for a-plane GaN and a-plane InN [43]. The Bragg spots appear with weak streaky lines in the RHEED patterns observed for a-plane GaN that confirms the reasonable smooth surface [44]. Spotty nature of a-plane InN RHEED pattern indicates the 3D growth of

#### 2.2. Nonpolar InN/GaN heterostructures

Growth of nonpolar a-plane InN/GaN heterostructures has been an important subject recently due to its potential improvement on the efficiency of III-nitride-based optoelectronic devices [32, 33]. However, growth of high-quality nonpolar InN/GaN heterostructures is challenging due to the low thermal decomposition temperature of InN film and high equilibrium vapor pressure of nitrogen. Despite the growth of high-quality nonpolar GaN films [23, 34, 35],

to grow active InN epilayers. The nitrogen flow rate and the forward RF power of the plasma

Figure 3(a) shows a 2θ-ω scan of InN films grown at different growth temperatures [29]. The peaks at 2θ = 31.3 and 65.5 are attributed to the (0002) and (0004) planes of the InN epilayers, whereas the peaks at 2θ = 34.56 and 72.81 are attributed to the (0002) and (0004) planes of the GaN templates. The peak at 2θ = 41.69 corresponds to the (0006) plane of sapphire substrate. The InN films grown at low (450C) and high (530C) temperatures show a peak at around 2θ = 33 indicating the presence of In metal. The presence of In metal could be due to the low migration of In at low temperature and more dissociation of InN at high temperature. The structural quality of the films was evaluated from the full width at half maximum (FWHM) of (0002) InN X-ray rocking curve (XRC). The rocking curves of the (0002) InN reflection are shown in Figure 3(b) [29]. Growth of InN epilayers at high temperature improved crystal quality, i.e., the FWHM of the (0002) InN XRC decreased from 532.8 to 450 arcsec corresponding to growth temperatures 450 and 500C. The low migration velocities of In and N atoms at low growth temperature are the most probable reason for the relatively inferior crystal quality [30]. On the other hand, growth of InN at higher temperature (530C) resulted in pronounced dissociation of InN, thus leading to high FWHM of 716.4 arcsec. The screw dislocation density of InN films, as calculated from the FWHM of the rocking curves, was

source were set to 0.5 sccm and 350 W. The indium BEP was 1.53 <sup>10</sup><sup>7</sup> mbar.

, 1.62 108

respectively. In addition, InN films show n-type conductivity with carrier concentrations in the

low growth temperature may be due to relatively poor crystallinity. On the other hand, the carrier concentration increased to 1.10 1019 cm<sup>3</sup> at high growth temperature of 530C,

The room temperature optical absorption spectra squared as a function of growth temperature is shown in Figure 3(c) [29]. The absorption spectrum exhibits characteristic interference fringes due to the underlying thicker GaN epilayer. The absorption edge of the InN films was estimated by extrapolating the linear part of the squared absorption down to the photon energy axis. The energy corresponding to this absorption edge is the amount of energy needed by an electron to make a vertical transition from the upper valence band to the Fermi surface in the conduction band. Therefore, this energy can be considered as the Fermi-level energy in the conduction band. A strong shift was observed in the absorption edge with a change in carrier concentration. This is usually referred as Burstein-Moss shift which is a commonly observed phenomenon in narrow-gap semiconductors owing to non-parabolic conduction band [31].

Growth of nonpolar a-plane InN/GaN heterostructures has been an important subject recently due to its potential improvement on the efficiency of III-nitride-based optoelectronic devices [32, 33]. However, growth of high-quality nonpolar InN/GaN heterostructures is challenging due to the low thermal decomposition temperature of InN film and high equilibrium vapor pressure of nitrogen. Despite the growth of high-quality nonpolar GaN films [23, 34, 35],

, and 4.12 108 cm<sup>2</sup> for samples A, B, C, and D,

. The high carrier concentration at

. As growth temperature increases from 450 to 500C, the carrier

found to be 2.27 108

176 Epitaxy

order of 1018 to 1019 cm<sup>3</sup>

, 2.03 108

concentration decreases from 1.57 1019 to 3.10 1018 cm<sup>3</sup>

which may be associated with the large dissociation of InN.

2.2. Nonpolar InN/GaN heterostructures

Figure 3. (a) 2θ-ω HRXRD scanning curve of InN films grown at different growth temperatures, (b) the XRC of the (0002) reflections of InN films grown at different growth temperatures, and (c) optical absorption spectra of InN films. Reprinted with permission from [29].

there are very few reports on the growth of nonpolar InN films [36, 37]. The earlier reports on the InN on r-sapphire substrates indicate the growth of cubic (001) [38, 39] and polar (0001) InN [40]. The nonpolar a-plane InN was demonstrated by using GaN buffer layer on r-plane sapphire [41, 42]. In this section, nonpolar (1 1 2 0) a-InN/GaN heterostructures were grown on rplane and m-plane sapphire substrates, respectively, using MBE system. Prior to the growth, thermal cleaning of the sapphire substrate was carried out at 850C inside MBE chamber for 30 min under ultrahigh vacuum. RF power and flow rate were kept constant at 350 W and 0.5 sccm, respectively. (1 1 2 0) a-GaN buffer layer was grown at 760C temperature. The growth of InN epilayers was carried out using two-step growth processes: growth of low-temperature InN buffer layer (~20 nm) on GaN under layer at 400C and growth of InN epilayers at different temperatures. The growth temperature of (1 1 2 0) a-InN varied from 470 to 530C.

Figure 4(a) show 2θ-ω scans of nonpolar (1 1 2 0) a-plane InN/GaN heterostructures grown on r-plane sapphire substrate at different growth temperatures [42]. The peaks at 2θ = 57.7 and 51.61 are assigned to the a-plane GaN and a-plane InN reflections, respectively, along with that from the r-plane sapphire substrate. The sample grown at 530C shows the peak at 32.9 which corresponds to In (101) reflections. Figure 4(b) shows the RHEED patterns in the azimuths [0001] for a-plane GaN and a-plane InN [43]. The Bragg spots appear with weak streaky lines in the RHEED patterns observed for a-plane GaN that confirms the reasonable smooth surface [44]. Spotty nature of a-plane InN RHEED pattern indicates the 3D growth of

2.3. Polar InGan/GaN heterostructures

performance of the InGaN-based devices [55, 56].

grown at 550C with In/Ga ~0.61.

Indium gallium nitride (InGaN), a ternary compound of III-nitride semiconductors, has received considerable attention due to its potential applications in optoelectronic devices [48–51]. The choice of InGaN as an active layer in high-performance optoelectronic devices is due to the advantage one gets in tuning the energy bandgap from visible to near-ultraviolet region by changing the In composition. The most challenging issues in InGaN-based nitride semiconductors include the spatial fluctuation of indium composition and the generation of dislocations at the interface of InGaN-based heterostructures due to the limited solubility of indium atom into GaN because of the difference in the In-N and Ga-N bond length [52–54]. Because of dislocations the non-radiative recombination increases, leading to the rapid decrease in the

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InGaN films of thicknesses around 200 nm were grown on 4 μm GaN/Al2O3 (0 0 0 1) templates using PAMBE. After the templates were chemically degreased, they were outgassed at 700C for 5 min in the presence of nitrogen plasma. Followed by thermal cleaning, the InGaN films were grown in a single step. The indium composition was varied by changing the In/Ga flux ratio and substrate temperature, whereas the nitrogen plasma conditions were constant, flow rate ~0.5 SCCM and forward RF power of 350 W. Figure 5(a) shows a 2θ-ω scan of as-grown InGaN films [57]. Well-resolved peaks corresponding to InGaN (0002) reflection was observed along with the GaN and Al2O3 peaks in all the samples. The peaks at 2θ = 34.56 and 41.69 attribute to the (0002) planes of the GaN films and (0006) plane of the sapphire substrate, respectively. The indium concentration in the as-grown InGaN films was determined by linearly interpolating the peak position of (0002) plane from their end binaries, by assuming that Vegard's law is valid [58, 59]. The indium concentration obtained for InGaN films with different growth conditions is given in Table 1. For a constant In/Ga flux ratio (0.61), the decrease in growth temperature 600 to 550C leads to the suppression of spinoidal decomposition. The transformation of multiple peaks to a single peak corresponding to InGaN (0002) planes is the evidence. The samples grown at high temperature, ~560C, with high In/Ga flux ratio (0.99) showed similar single peaks. However, the presence of In (101) peak states that metallic In was also present. Further, it was also observed that at growth temperature ~560C, the increase in In/Ga ratio from 0.61 to 0.99 not only led to increase in the In incorporation but also decreased the spinoidal decomposition. Therefore, we could obtain high-quality, metalfree, single composition InGaN epilayers with 23% indium (sample E) on GaN/c-Al2O3 when

The X-ray rocking curve (XRC) was carried out to see the structural quality of the InGaN film. Figure 5(b) shows the XRC of the (0002) InGaN reflection of sample E [57], and the corresponding FWHM was found to be 390 arcsec, indicating the high quality of the as-grown InGaN film. This value is comparable to the values in literature [60, 61]. Though this value indicates the quality of bulk film, it doesn't tell us much about the interface dislocations. The influence of such interface defects on the barrier height, and the ideality factor will be discussed in further sections. Figure 5(c) shows the room temperature PL spectra of sample E [57]. As we can see from the PL spectra, the emission peak at 2.48 eV corresponds to the free excitonic transition between valence and conduction bands of InGaN film. Also, one can notice the presence of fringes around this peak, which are a result of Fabry-Perot interference, hence

Figure 4. (a) 2θ-ω HRXRD scans of nonpolar a-plane InN film grown on GaN/r-Al2O3 substrate at different temperatures, (b) RHEED patterns of a-plane (1120) GaN and a-plane (1120) InN taken along (0 0 0 1) azimuth, (c) rocking curve FWHM of nonpolar a-plane InN grown at 500C with different azimuth angles, and (d) optical absorption spectra of nonpolar a-plane InN. Reprinted with permission from [42, 43].

InN [43]. The X-ray rocking was carried out to see the structural quality of the samples. The FWHM of the rocking curve along with different azimuth angles has been plotted in Figure 4(c) for InN film grown at 500C [42]. The variation of FWHM was found to be M-type with respect to the azimuth angles. The FWHM of the (1 1 2 0) reflection was found to be strongly dependent on the azimuth angle with respect to the scattering plane. Azimuth angle was defined as zero when the incident beam is parallel to the [0001] direction. Darakchieva et al. reported similar kind of M-type behavior in a-plane InN epilayers grown on r-sapphire substrates with GaN buffer layer [45]. The FWHM of (1 1 2 0) reflection along [0 0 0 1] and [1 1 0 0] directions was determined, and the values are 0.67 and 0.85, respectively.

Absorption spectroscopy was used to determine the bandgap of nonpolar a-plane InN films and is shown in Figure 4(d) [42]. The direct optical bandgap for InN can be investigated by fitting the absorption data. The bandgap of InN epilayers grown at 470, 500, and 530C as obtained by fitting the absorption data is found to be 0.81, 0.74, and 0.78 eV, respectively. This shows that the bandgap of the samples grown at low temperatures is blue shifted with respect to the bulk, which could be due to the high background carrier concentration [46]. The carrier concentration of InN films was estimated by using Hall measurements, and the carrier concentrations was found to be in the order of 1018 to 1019 cm<sup>3</sup> with growth temperature. The film grown at 500C showed the lowest carrier concentration, whereas the film grown at 470C showed a higher carrier concentration due to the poor crystallinity. The carrier concentration in the film grown at 530C was also found to be higher and could be as a result of high dissociation rate of InN at that temperature [47].

#### 2.3. Polar InGan/GaN heterostructures

InN [43]. The X-ray rocking was carried out to see the structural quality of the samples. The FWHM of the rocking curve along with different azimuth angles has been plotted in Figure 4(c) for InN film grown at 500C [42]. The variation of FWHM was found to be M-type with respect to the azimuth angles. The FWHM of the (1 1 2 0) reflection was found to be strongly dependent on the azimuth angle with respect to the scattering plane. Azimuth angle was defined as zero when the incident beam is parallel to the [0001] direction. Darakchieva et al. reported similar kind of M-type behavior in a-plane InN epilayers grown on r-sapphire substrates with GaN buffer layer [45]. The FWHM of (1 1 2 0) reflection along [0 0 0 1] and [1 1

Figure 4. (a) 2θ-ω HRXRD scans of nonpolar a-plane InN film grown on GaN/r-Al2O3 substrate at different temperatures, (b) RHEED patterns of a-plane (1120) GaN and a-plane (1120) InN taken along (0 0 0 1) azimuth, (c) rocking curve FWHM of nonpolar a-plane InN grown at 500C with different azimuth angles, and (d) optical absorption spectra of

Absorption spectroscopy was used to determine the bandgap of nonpolar a-plane InN films and is shown in Figure 4(d) [42]. The direct optical bandgap for InN can be investigated by fitting the absorption data. The bandgap of InN epilayers grown at 470, 500, and 530C as obtained by fitting the absorption data is found to be 0.81, 0.74, and 0.78 eV, respectively. This shows that the bandgap of the samples grown at low temperatures is blue shifted with respect to the bulk, which could be due to the high background carrier concentration [46]. The carrier concentration of InN films was estimated by using Hall measurements, and the carrier concentrations was found to be in the order of 1018 to 1019 cm<sup>3</sup> with growth temperature. The film grown at 500C showed the lowest carrier concentration, whereas the film grown at 470C showed a higher carrier concentration due to the poor crystallinity. The carrier concentration in the film grown at 530C was also found to be higher and could be as a result of high

0 0] directions was determined, and the values are 0.67 and 0.85, respectively.

dissociation rate of InN at that temperature [47].

nonpolar a-plane InN. Reprinted with permission from [42, 43].

178 Epitaxy

Indium gallium nitride (InGaN), a ternary compound of III-nitride semiconductors, has received considerable attention due to its potential applications in optoelectronic devices [48–51]. The choice of InGaN as an active layer in high-performance optoelectronic devices is due to the advantage one gets in tuning the energy bandgap from visible to near-ultraviolet region by changing the In composition. The most challenging issues in InGaN-based nitride semiconductors include the spatial fluctuation of indium composition and the generation of dislocations at the interface of InGaN-based heterostructures due to the limited solubility of indium atom into GaN because of the difference in the In-N and Ga-N bond length [52–54]. Because of dislocations the non-radiative recombination increases, leading to the rapid decrease in the performance of the InGaN-based devices [55, 56].

InGaN films of thicknesses around 200 nm were grown on 4 μm GaN/Al2O3 (0 0 0 1) templates using PAMBE. After the templates were chemically degreased, they were outgassed at 700C for 5 min in the presence of nitrogen plasma. Followed by thermal cleaning, the InGaN films were grown in a single step. The indium composition was varied by changing the In/Ga flux ratio and substrate temperature, whereas the nitrogen plasma conditions were constant, flow rate ~0.5 SCCM and forward RF power of 350 W. Figure 5(a) shows a 2θ-ω scan of as-grown InGaN films [57]. Well-resolved peaks corresponding to InGaN (0002) reflection was observed along with the GaN and Al2O3 peaks in all the samples. The peaks at 2θ = 34.56 and 41.69 attribute to the (0002) planes of the GaN films and (0006) plane of the sapphire substrate, respectively. The indium concentration in the as-grown InGaN films was determined by linearly interpolating the peak position of (0002) plane from their end binaries, by assuming that Vegard's law is valid [58, 59]. The indium concentration obtained for InGaN films with different growth conditions is given in Table 1. For a constant In/Ga flux ratio (0.61), the decrease in growth temperature 600 to 550C leads to the suppression of spinoidal decomposition. The transformation of multiple peaks to a single peak corresponding to InGaN (0002) planes is the evidence. The samples grown at high temperature, ~560C, with high In/Ga flux ratio (0.99) showed similar single peaks. However, the presence of In (101) peak states that metallic In was also present. Further, it was also observed that at growth temperature ~560C, the increase in In/Ga ratio from 0.61 to 0.99 not only led to increase in the In incorporation but also decreased the spinoidal decomposition. Therefore, we could obtain high-quality, metalfree, single composition InGaN epilayers with 23% indium (sample E) on GaN/c-Al2O3 when grown at 550C with In/Ga ~0.61.

The X-ray rocking curve (XRC) was carried out to see the structural quality of the InGaN film. Figure 5(b) shows the XRC of the (0002) InGaN reflection of sample E [57], and the corresponding FWHM was found to be 390 arcsec, indicating the high quality of the as-grown InGaN film. This value is comparable to the values in literature [60, 61]. Though this value indicates the quality of bulk film, it doesn't tell us much about the interface dislocations. The influence of such interface defects on the barrier height, and the ideality factor will be discussed in further sections. Figure 5(c) shows the room temperature PL spectra of sample E [57]. As we can see from the PL spectra, the emission peak at 2.48 eV corresponds to the free excitonic transition between valence and conduction bands of InGaN film. Also, one can notice the presence of fringes around this peak, which are a result of Fabry-Perot interference, hence

layer was calculated using Vegard's law and was found to be 22% in sample E, which is in correlation with the value as estimated by HRXRD. Figure 5(d) is the high-resolution transmission electron microscopy (HRTEM), which showing an extremely sharp interface and the growth plane (0001) is perpendicular to the growth direction. High-quality interface as seen from the HRTEM image justifies the less defect densities and high crystallinity of the InGaN films.

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InGaN alloys of various compositions are being optimized as materials for the fabrication of light-emitting diodes which are active in the entire visible spectrum extending up to ultraviolet wavelengths [32, 64]. Indium-rich InGaN alloys are now being considered potential candidates for longer wavelength emitters, thermionic emitters, multi-junction solar cells, etc. [65–67]. A concern with polar heterostructures is the intrinsic and strong polarization fields resident in the lattice. To overcome such polarization effect, substrates oriented in nonpolar directions, i.e., (1 0 1 0) m-plane or (1 1 2 0) a-plane, are used. Devices grown on these orientations are receiving a lot of focus due to this enhanced behavior. However, a slight compromise in terms of quality has to be made because of the large mismatch in lattice constants, thermal expansion coefficients, and elastic constants of InGaN and GaN. In addition, the large lattice constant mismatch between GaN and InN (~11%) results in a phase separation in InGaN alloys which has been theoretically predicted and experimentally observed [68]. This makes the growth of InGaN very challenging, especially for higher concentration of indium (>20%) [58]. Growth parameters such as growth temperature, growth rate, and flux ratio are seen to drastically affect the indium incorporation in InGaN films. During growth of the InGaN alloys, the evaporation of indium species from the surface will be suppressed at lower temperatures and leads to higher growth rates as the indium species become trapped by the growing layer [69, 70]. As the indium content in the InGaN alloy increases, the material quality degrades due to phase separation, inhomogeneity of solid solution, and indium metal droplets. Few reports are available in the literature for the study of the nonpolar InGaN-based multiple quantum well (MQW) structures and other devices. Song et al. studied the effect of periodicity of a-plane InGaN/GaN multiple quantum wells on the output power of the LEDs

In this section, InxGa1xN films of thicknesses around 125 nm were grown on 125 nm a-GaN (1120)/r-Al2O3 (1102) by PAMBE system. A two-step growth process was employed to grow a-GaN film, which constituted the growth of a 20 nm thin low-temperature GaN buffer layer at 500C and a ~125nm GaN epilayer at 760C. The nitrogen plasma RF power and N2 flow rate were kept constant at 350 W and 1 sccm, respectively, throughout the growth duration. The gallium beam equivalent pressure (BEP) was kept at 5.6 <sup>10</sup><sup>7</sup> mbar, corresponding to the growth in the slightly nitrogen-rich region. InxGa1xN films were grown on top of the a-GaN layer at the growth conditions as tabulated in Table 2. Three InxGa1xN films (samples A, B, and C) with different indium compositions were grown by varying the In/ Ga flux ratios and growth temperatures. During the InxGa1-xN growth, the nitrogen flow rate

and RF power of the plasma source were kept at 0.5 sccm and 350 W, respectively.

Figure 6(a) shows the 2θ-ω HRXRD scan of the grown films, which confirmed the growth of nonpolar a-InGaN epilayer oriented in the (1 1 2 0) direction and a-GaN epilayer oriented in

2.4. Nonpolar InGaN/GaN heterostructures

grown by MOCVD [71].

Figure 5. (a) 2θ-ω HRXRD scanning curve of InGaN films grown on GaN, (b) X-ray rocking curve of the (0002) plane of InGaN film (sample-E), (c) the room temperature photoluminescence spectra of InGaN film grown on GaN (sample E), and (d) HRTEM of InGaN/GaN interface. Reprinted with permission from [57].


Table 1. Growth parameters used for the growth of InGaN films.

suggesting a smooth and abrupt interface. The film also exhibits a peak at 3.41 eV, which corresponds to underlying GaN layer. A second peak at 2.94 eV might correspond to the thin InGaN strain-relieving layer with ~6% of In [62, 63]. The indium composition of the InGaN layer was calculated using Vegard's law and was found to be 22% in sample E, which is in correlation with the value as estimated by HRXRD. Figure 5(d) is the high-resolution transmission electron microscopy (HRTEM), which showing an extremely sharp interface and the growth plane (0001) is perpendicular to the growth direction. High-quality interface as seen from the HRTEM image justifies the less defect densities and high crystallinity of the InGaN films.

#### 2.4. Nonpolar InGaN/GaN heterostructures

suggesting a smooth and abrupt interface. The film also exhibits a peak at 3.41 eV, which corresponds to underlying GaN layer. A second peak at 2.94 eV might correspond to the thin InGaN strain-relieving layer with ~6% of In [62, 63]. The indium composition of the InGaN

Figure 5. (a) 2θ-ω HRXRD scanning curve of InGaN films grown on GaN, (b) X-ray rocking curve of the (0002) plane of InGaN film (sample-E), (c) the room temperature photoluminescence spectra of InGaN film grown on GaN (sample E),

Sample ID In/Ga flux ratios Growth temperature (C) In composition Sample A 0.61 600 11, 16, 24% Sample B 0.99 540 28, 35% Sample C 0.61 560 20, 24% Sample D 0.99 560 33% (In) Sample E 0.61 550 23%

and (d) HRTEM of InGaN/GaN interface. Reprinted with permission from [57].

180 Epitaxy

Table 1. Growth parameters used for the growth of InGaN films.

InGaN alloys of various compositions are being optimized as materials for the fabrication of light-emitting diodes which are active in the entire visible spectrum extending up to ultraviolet wavelengths [32, 64]. Indium-rich InGaN alloys are now being considered potential candidates for longer wavelength emitters, thermionic emitters, multi-junction solar cells, etc. [65–67]. A concern with polar heterostructures is the intrinsic and strong polarization fields resident in the lattice. To overcome such polarization effect, substrates oriented in nonpolar directions, i.e., (1 0 1 0) m-plane or (1 1 2 0) a-plane, are used. Devices grown on these orientations are receiving a lot of focus due to this enhanced behavior. However, a slight compromise in terms of quality has to be made because of the large mismatch in lattice constants, thermal expansion coefficients, and elastic constants of InGaN and GaN. In addition, the large lattice constant mismatch between GaN and InN (~11%) results in a phase separation in InGaN alloys which has been theoretically predicted and experimentally observed [68]. This makes the growth of InGaN very challenging, especially for higher concentration of indium (>20%) [58]. Growth parameters such as growth temperature, growth rate, and flux ratio are seen to drastically affect the indium incorporation in InGaN films. During growth of the InGaN alloys, the evaporation of indium species from the surface will be suppressed at lower temperatures and leads to higher growth rates as the indium species become trapped by the growing layer [69, 70]. As the indium content in the InGaN alloy increases, the material quality degrades due to phase separation, inhomogeneity of solid solution, and indium metal droplets. Few reports are available in the literature for the study of the nonpolar InGaN-based multiple quantum well (MQW) structures and other devices. Song et al. studied the effect of periodicity of a-plane InGaN/GaN multiple quantum wells on the output power of the LEDs grown by MOCVD [71].

In this section, InxGa1xN films of thicknesses around 125 nm were grown on 125 nm a-GaN (1120)/r-Al2O3 (1102) by PAMBE system. A two-step growth process was employed to grow a-GaN film, which constituted the growth of a 20 nm thin low-temperature GaN buffer layer at 500C and a ~125nm GaN epilayer at 760C. The nitrogen plasma RF power and N2 flow rate were kept constant at 350 W and 1 sccm, respectively, throughout the growth duration. The gallium beam equivalent pressure (BEP) was kept at 5.6 <sup>10</sup><sup>7</sup> mbar, corresponding to the growth in the slightly nitrogen-rich region. InxGa1xN films were grown on top of the a-GaN layer at the growth conditions as tabulated in Table 2. Three InxGa1xN films (samples A, B, and C) with different indium compositions were grown by varying the In/ Ga flux ratios and growth temperatures. During the InxGa1-xN growth, the nitrogen flow rate and RF power of the plasma source were kept at 0.5 sccm and 350 W, respectively.

Figure 6(a) shows the 2θ-ω HRXRD scan of the grown films, which confirmed the growth of nonpolar a-InGaN epilayer oriented in the (1 1 2 0) direction and a-GaN epilayer oriented in


Table 2. Growth conditions for a-InxGa1xN/a-GaN/r-sapphire substrate.

the (1 1 2 0) direction on (1 1 0 2) r-plane sapphire. The peak at 2θ = 56.64 for sample A, 2θ = 56.59 for sample B, and 2θ = 56.36 for sample C was assigned to (1 1 2 0) InxGa1-xN which corresponds to the different composition of indium in the grown films. The calculated compositions of the samples were In0.17Ga0.83N (sample A), In0.19Ga0.81N (sample B), and In0.22Ga0.78N (sample C). We observe that by decreasing the indium flux during the growth from 1.17 <sup>10</sup><sup>7</sup> to 8.87 <sup>10</sup><sup>8</sup> , we could incorporate more indium into the InxGa1xN alloy. Higher indium incorporation was also observed when the substrate temperature was decreased from 550C to 540C keeping the indium flux constant. Growth parameters play a very critical role in controlling the composition of the InxGa1xN films. For the given set of growth parameters, we have not observed any phase separation or indium segregation in the grown InxGa1xN films.

X-ray rocking curve (ω) analysis was carried out to determine the crystal quality of the as-grown structures. Rocking curves along different azimuth angles have been recorded, and RC FWHM versus azimuth angle has been plotted for all the samples as shown in Figure 6(b). The RC FWHM of the reflection along the normal varied with the azimuth angle and showed an M-type behavior [72]. The measured FWHM value of (1 1 2 0) reflection of In0.22Ga0.78N along [0001] direction defined as azimuth angle 0and along [1100] direction defined as azimuth angle 90 was found to be 0.532 and 0.703, respectively. The RC FWHM values of (1 1 2 0) GaN reflections were found to be 0.47 and 0.52 along azimuth angle 0 corresponding to reflection along [0001] direction and 90 corresponding to reflection along [1100] direction, respectively. The reason behind the broadening of InGaN rocking curves could be attributed to the presence of defects such as partial dislocations and stacking faults, thus suggesting that the crystalline quality of nonpolar InGaN films is reduced with the increase in indium incorporation [73]. Crosssectional plan view of TEM image obtained in bright field is shown in Figure 6(c), which shows a clear interface of the a-InGaN/a-GaN/r-sapphire structure. The thickness of the each layer grown is confirmed from the image to be around 125 nm each. Basal stacking faults (BSF) are visible as thin lines, and they arise from the low-temperature nucleation layer. Due to the anisotropic nature of the growing surface, the nonpolar GaN typically has a high density of BSFs. The room temperature PL spectrum of InxGa1xN films is shown in Figure 6(d). The position of the near-band-edge emission (NBE) of the InxGa1-xN as observed for the three samples were 2.67, 2.59, and 2.56 eV for samples A, B, and C, respectively. Using Vegard's law, the values of indium fraction were found to be 0.20, 0.22, and 0.23 for samples A, B, and C, respectively.

in the reciprocal space if additional mosaic tilt exists in the sample. From Figure 7(a)–(c), the RLPs along the symmetric (1 1 2 0) reflection of all samples are all broadened in the Qx direction with negligible inclination, indicating that the dominant broadening mechanism for it is the limited mosaic block size. Asymmetric reflection are expected to be elongated parallel to the lateral scattering vector if the material experiences broadening due to short-sized mosaic blocks and could be also inclined if an additional mosaicity tilting exists in the material [75]. To understand the strain in the film with respect to the substrate, asymmetric RSM along (1 0 1 0) a-GaN scans were obtained and are shown in Figure 7(d)–(f). Shift of the center of each peak along the Qz axis in an asymmetric scan gives a direct evidence of the strain present between the layers. Splitting of the GaN peak for all compositions implies a formation of a thin layer of high gallium composition InxGa1xN layer which is highly strained with respect to the GaN

InxGa1xN films.

Figure 6. (a) HRXRD 2θ-ω scan of nonpolar (1120) a-InGaN epilayer grown on (1120) a-GaN/(1102) r-plane sapphire, (b) FWHM of the rocking curve along with variation in azimuth angles for InGaN and GaN layer, (c) bright-field plan-view cross-sectional image of a-InGaN/a-GaN/r-sapphire substrate, and (d) room temperature PL spectrum of

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Reciprocal space maps (RSMs) were recorded to look into the microstructure and strain present in the films. The reciprocal lattice points (RLPs) will be elongated along the Qx axis if the broadening is caused by the limited mosaic block dimensions [74]. The RLPs will be tilted

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the (1 1 2 0) direction on (1 1 0 2) r-plane sapphire. The peak at 2θ = 56.64 for sample A, 2θ = 56.59 for sample B, and 2θ = 56.36 for sample C was assigned to (1 1 2 0) InxGa1-xN which corresponds to the different composition of indium in the grown films. The calculated compositions of the samples were In0.17Ga0.83N (sample A), In0.19Ga0.81N (sample B), and In0.22Ga0.78N (sample C). We observe that by decreasing the indium flux during the growth

Sample ID Gallium BEP (mbar) Indium BEP (mbar) Growth temp. (C)

(A) 1.23 <sup>10</sup><sup>7</sup> 1.17 <sup>10</sup><sup>7</sup> <sup>550</sup> (B) 1.23 <sup>10</sup><sup>7</sup> 8.87 <sup>10</sup><sup>8</sup> <sup>550</sup> (C) 1.23 <sup>10</sup><sup>7</sup> 8.87 <sup>10</sup><sup>8</sup> <sup>540</sup>

Table 2. Growth conditions for a-InxGa1xN/a-GaN/r-sapphire substrate.

Higher indium incorporation was also observed when the substrate temperature was decreased from 550C to 540C keeping the indium flux constant. Growth parameters play a very critical role in controlling the composition of the InxGa1xN films. For the given set of growth parameters, we have not observed any phase separation or indium segregation in the

X-ray rocking curve (ω) analysis was carried out to determine the crystal quality of the as-grown structures. Rocking curves along different azimuth angles have been recorded, and RC FWHM versus azimuth angle has been plotted for all the samples as shown in Figure 6(b). The RC FWHM of the reflection along the normal varied with the azimuth angle and showed an M-type behavior [72]. The measured FWHM value of (1 1 2 0) reflection of In0.22Ga0.78N along [0001] direction defined as azimuth angle 0and along [1100] direction defined as azimuth angle 90 was found to be 0.532 and 0.703, respectively. The RC FWHM values of (1 1 2 0) GaN reflections were found to be 0.47 and 0.52 along azimuth angle 0 corresponding to reflection along [0001] direction and 90 corresponding to reflection along [1100] direction, respectively. The reason behind the broadening of InGaN rocking curves could be attributed to the presence of defects such as partial dislocations and stacking faults, thus suggesting that the crystalline quality of nonpolar InGaN films is reduced with the increase in indium incorporation [73]. Crosssectional plan view of TEM image obtained in bright field is shown in Figure 6(c), which shows a clear interface of the a-InGaN/a-GaN/r-sapphire structure. The thickness of the each layer grown is confirmed from the image to be around 125 nm each. Basal stacking faults (BSF) are visible as thin lines, and they arise from the low-temperature nucleation layer. Due to the anisotropic nature of the growing surface, the nonpolar GaN typically has a high density of BSFs. The room temperature PL spectrum of InxGa1xN films is shown in Figure 6(d). The position of the near-band-edge emission (NBE) of the InxGa1-xN as observed for the three samples were 2.67, 2.59, and 2.56 eV for samples A, B, and C, respectively. Using Vegard's law, the values of indium

fraction were found to be 0.20, 0.22, and 0.23 for samples A, B, and C, respectively.

Reciprocal space maps (RSMs) were recorded to look into the microstructure and strain present in the films. The reciprocal lattice points (RLPs) will be elongated along the Qx axis if the broadening is caused by the limited mosaic block dimensions [74]. The RLPs will be tilted

, we could incorporate more indium into the InxGa1xN alloy.

from 1.17 <sup>10</sup><sup>7</sup> to 8.87 <sup>10</sup><sup>8</sup>

grown InxGa1xN films.

182 Epitaxy

Figure 6. (a) HRXRD 2θ-ω scan of nonpolar (1120) a-InGaN epilayer grown on (1120) a-GaN/(1102) r-plane sapphire, (b) FWHM of the rocking curve along with variation in azimuth angles for InGaN and GaN layer, (c) bright-field plan-view cross-sectional image of a-InGaN/a-GaN/r-sapphire substrate, and (d) room temperature PL spectrum of InxGa1xN films.

in the reciprocal space if additional mosaic tilt exists in the sample. From Figure 7(a)–(c), the RLPs along the symmetric (1 1 2 0) reflection of all samples are all broadened in the Qx direction with negligible inclination, indicating that the dominant broadening mechanism for it is the limited mosaic block size. Asymmetric reflection are expected to be elongated parallel to the lateral scattering vector if the material experiences broadening due to short-sized mosaic blocks and could be also inclined if an additional mosaicity tilting exists in the material [75]. To understand the strain in the film with respect to the substrate, asymmetric RSM along (1 0 1 0) a-GaN scans were obtained and are shown in Figure 7(d)–(f). Shift of the center of each peak along the Qz axis in an asymmetric scan gives a direct evidence of the strain present between the layers. Splitting of the GaN peak for all compositions implies a formation of a thin layer of high gallium composition InxGa1xN layer which is highly strained with respect to the GaN

2.5. InGaN/Si heterostructures

tions (TDs) [73].

n-n heterojunctions have found use in different applications such as photodetectors, lightemitting diodes, solar cells, injection lasers, etc. [76–80]. So far the most extensively studied material systems are ZnO/Si, InxGa1-xAs, InxGa1-xSb, Ge-Si, etc. [81, 82]. Very few groups have reported studies on isotype heterojunctions of InxGa1xN system [83, 84]. The large lattice mismatch of InGaN with silicon substrates has hindered the realization of silicon-based commercial devices and is limited to sapphire which happens to be an insulator. Therefore, obtaining smooth and abrupt heterojunctions with minimum density of interface defects by overcoming the lattice mismatch issue has been a concern of great interest for many researchers across the globe. Several attempts have been made to demonstrate low power consuming or self-powered photodetectors [85, 86]. There are a few reports on the UV photodetection using InGaN as an active layer [87]. In this work, we report the development of a UV photodetector which is operated at zero bias. The device comprises of a simple n-InGaN/n-Si heterojunction. The role of interface defects originating due to the large lattice mismatch, such as traps, in modulating the built-in electric field driven photoresponse has been discussed. The n-Si (111) substrates of 1 1 cm2 in size were cleaned chemically by trichloroethylene, acetone, and methanol and were dipped in 5% HF for 60 s to remove the native oxide prior to loading in the growth chamber. Thermal cleaning of Si (111) was carried out at 900C for the removal of native oxide layer. The substrate temperature was further reduced to 550C, and growth was carried out for 2 h without any intermediate steps. The indium (In) beam equivalent pressure (BEP) and gallium (Ga) BEP, nitrogen flow, and plasma power were kept at 8.52 <sup>10</sup><sup>8</sup> mbar, 1.2 <sup>10</sup><sup>7</sup> mbar, 1 sccm, and 350 W, respectively. Various experimental techniques were employed to study the as-grown samples as mentioned in upcoming sections. Circular aluminum contacts with diameters of 600 μm were then deposited by thermal evaporation on the InGaN films and Si (111) substrate with the help of a

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physical mask to study the current-voltage and photoresponse studies.

Figure 8(a) shows the 2θ-ω HRXRD scan of InGaN epilayers on Si (111) substrates [88]. The peaks at 2θ = 28.45 and 58.86 attribute to Si (111) and Si (222) reflections, respectively. The peaks at 2θ = 34.05 and 71.89 correspond to (0002) and (0004) reflections of InGaN. Also, no secondary phases or residual indium or InN phases were found, confirming that the films are single crystalline. The In composition was determined to be 15% from the linear interpolation of the 2θ peak positions of (0002) GaN (34.59) and (0002) InN (31.22). In Figure 8(b) the rocking curve FWHM value (degree) of asymmetric reflection (2.16) is higher than that of the symmetric reflection (2.05), which attributes to the presence of large edge threading disloca-

The room temperature PL spectra of the InxGa1xN films are shown in Figure 8(c). Near-band-edge emission peaks are observed at 502.12 nm corresponding to the bandgap of 2.46 eV. Due to stoke shift of the PL spectra and discrepancies in the exact value of InN, bandgap [89] further adds ambiguity to the exact In content determination from PL spectra. A small hump is observed around 425 nm which might be due to the initial layers with a large number of dislocations arising due to the large lattice mismatch of Si and InGaN films [90]. One can also correlate the presence of a large number of defects from the FWHM values of both symmetric

Figure 7. (a)–(c) RSM of symmetric (1 1 2 0) of InxGa1xN/GaN/r-sapphire and (d)–(f) RSM of asymmetric (1 0 1 0) InxGa1xN/GaN layers.

layer. Hence, from the asymmetric RSM scans, it is understood that after a thin layer of InxGa1xN acting as a buffer layer, the InxGa1xN film with In concentration of 17, 19 and 22% was formed in samples A, B, and C, respectively. As the In concentration is increased, the InGaN film is showing signs of relaxation which indirectly means increase in the dislocation density, reducing the crystallinity of the layer which is consistent with the observation of the variation of FWHM in rocking curve measurements.

#### 2.5. InGaN/Si heterostructures

layer. Hence, from the asymmetric RSM scans, it is understood that after a thin layer of InxGa1xN acting as a buffer layer, the InxGa1xN film with In concentration of 17, 19 and 22% was formed in samples A, B, and C, respectively. As the In concentration is increased, the InGaN film is showing signs of relaxation which indirectly means increase in the dislocation density, reducing the crystallinity of the layer which is consistent with the observation of the

Figure 7. (a)–(c) RSM of symmetric (1 1 2 0) of InxGa1xN/GaN/r-sapphire and (d)–(f) RSM of asymmetric (1 0 1 0)

variation of FWHM in rocking curve measurements.

InxGa1xN/GaN layers.

184 Epitaxy

n-n heterojunctions have found use in different applications such as photodetectors, lightemitting diodes, solar cells, injection lasers, etc. [76–80]. So far the most extensively studied material systems are ZnO/Si, InxGa1-xAs, InxGa1-xSb, Ge-Si, etc. [81, 82]. Very few groups have reported studies on isotype heterojunctions of InxGa1xN system [83, 84]. The large lattice mismatch of InGaN with silicon substrates has hindered the realization of silicon-based commercial devices and is limited to sapphire which happens to be an insulator. Therefore, obtaining smooth and abrupt heterojunctions with minimum density of interface defects by overcoming the lattice mismatch issue has been a concern of great interest for many researchers across the globe. Several attempts have been made to demonstrate low power consuming or self-powered photodetectors [85, 86]. There are a few reports on the UV photodetection using InGaN as an active layer [87]. In this work, we report the development of a UV photodetector which is operated at zero bias. The device comprises of a simple n-InGaN/n-Si heterojunction. The role of interface defects originating due to the large lattice mismatch, such as traps, in modulating the built-in electric field driven photoresponse has been discussed. The n-Si (111) substrates of 1 1 cm2 in size were cleaned chemically by trichloroethylene, acetone, and methanol and were dipped in 5% HF for 60 s to remove the native oxide prior to loading in the growth chamber. Thermal cleaning of Si (111) was carried out at 900C for the removal of native oxide layer. The substrate temperature was further reduced to 550C, and growth was carried out for 2 h without any intermediate steps. The indium (In) beam equivalent pressure (BEP) and gallium (Ga) BEP, nitrogen flow, and plasma power were kept at 8.52 <sup>10</sup><sup>8</sup> mbar, 1.2 <sup>10</sup><sup>7</sup> mbar, 1 sccm, and 350 W, respectively. Various experimental techniques were employed to study the as-grown samples as mentioned in upcoming sections. Circular aluminum contacts with diameters of 600 μm were then deposited by thermal evaporation on the InGaN films and Si (111) substrate with the help of a physical mask to study the current-voltage and photoresponse studies.

Figure 8(a) shows the 2θ-ω HRXRD scan of InGaN epilayers on Si (111) substrates [88]. The peaks at 2θ = 28.45 and 58.86 attribute to Si (111) and Si (222) reflections, respectively. The peaks at 2θ = 34.05 and 71.89 correspond to (0002) and (0004) reflections of InGaN. Also, no secondary phases or residual indium or InN phases were found, confirming that the films are single crystalline. The In composition was determined to be 15% from the linear interpolation of the 2θ peak positions of (0002) GaN (34.59) and (0002) InN (31.22). In Figure 8(b) the rocking curve FWHM value (degree) of asymmetric reflection (2.16) is higher than that of the symmetric reflection (2.05), which attributes to the presence of large edge threading dislocations (TDs) [73].

The room temperature PL spectra of the InxGa1xN films are shown in Figure 8(c). Near-band-edge emission peaks are observed at 502.12 nm corresponding to the bandgap of 2.46 eV. Due to stoke shift of the PL spectra and discrepancies in the exact value of InN, bandgap [89] further adds ambiguity to the exact In content determination from PL spectra. A small hump is observed around 425 nm which might be due to the initial layers with a large number of dislocations arising due to the large lattice mismatch of Si and InGaN films [90]. One can also correlate the presence of a large number of defects from the FWHM values of both symmetric

vital to study the transport properties of the semiconductor heterostructure-based Schottky junctions. The InN/GaN heterostructure has a large conduction band offset, thus ensuring effective blockage of the conduction current over the barriers [91, 92], which help for the development of electronic devices operating in THz frequency range [91]. Hence, studying the transport properties across the InN/GaN interface is very important from device view point. Chen et al. [93] studied temperature-dependent current-voltage characteristics of InN/GaN-based Schottky junctions in the range of 300–400K and found that the barrier height (~1.25 eV) and the ideality factor (~1.25) are nearly temperature independent. Similarly, Wang et al. [94] employed capacitance-voltage measurement technique to determine the Schottky barrier height to be 0.94 eV at room temperature. This section presents the study of the temperature-dependent electrical transport properties of InN/GaN heterostructures and observed the temperature-dependent barrier height and the ideality factor. Figure 9(a) shows the room temperature J-V(current density-voltage) characteristics for the junction. The junction between InN and GaN exhibits a rectifying behavior which suggests an existence of Schottky

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187

To investigate further, we have studied the temperature-dependent J-V (J-V-T) characteristics ranging from 200 to 500 K and is given in Figure 9(b) [95]. In the present case, the GaN film is

GaN interface, which is due to the tunneling of charge carriers through the interface. This kind of behavior reveals that the current transport is primarily dominated by thermionic field emission (TFE) mechanism, where the carriers tunnel from GaN to InN. According to the transport theory, TFE dominates only when E00/kT ≈ 1, where E00 is the characteristic tunneling energy which determines the nature of conduction mechanism. When E00/kT ≈ 1, the Boltzmann distribution tail of thermionic emission drops off by a factor of exp(1), which is much faster than the decrease rate of the tunneling probability. On the other hand, thermionic emission (TE) is predominant when E00/kT << 1 because the tunneling probability drops off faster than TE [97]. In the present case, the value of the barrier height and the ideality factor (η)

Figure 9. (a) The room temperature J–V characteristics of InN/GaN Schottky junction and (b) the forward J–V character-

istics with TFE fitting as a function of measurement temperature. Reprinted with permission from [95, 96].

) resulting in a lower barrier height at the InN/

barrier height at the junction [95, 96].

highly doped with silicon (ND ~ 1.4 1018 cm<sup>3</sup>

Figure 8. (a) HRXRD 2θ-ω scan of InGaN on bare Si (111), (b) X-ray rocking curve of (0002) and (1011) reflections, (c) photoluminescence spectra of InGaN/Si, and (d) cross-sectional TEM of InGaN/Si heterostructure. Reprinted with permission from [88].

and asymmetric XRCs. The contrast at the interface in the TEM image in Figure 8(d) clearly indicates the presence of initial low-contrast Ga-rich layers which possibly attribute to the ~425 nm peak in PL spectra. The thickness was estimated to be ~100 nm from the TEM image. The other dark regions are formed as a result from the Ga ion beam damage during sample thinning.

## 3. Transport properties of III-nitride heterostructure-based Schottky junctions

#### 3.1. Polar and nonpolar InN/GaN heterostructures.

The semiconductor heterostructure exhibits the Schottky barrier at the interface due to the formation of conduction band offset because of their different bandgap values. The concepts on the band offset are directly transferable to the Schottky barrier height problems. Hence, it is vital to study the transport properties of the semiconductor heterostructure-based Schottky junctions. The InN/GaN heterostructure has a large conduction band offset, thus ensuring effective blockage of the conduction current over the barriers [91, 92], which help for the development of electronic devices operating in THz frequency range [91]. Hence, studying the transport properties across the InN/GaN interface is very important from device view point. Chen et al. [93] studied temperature-dependent current-voltage characteristics of InN/GaN-based Schottky junctions in the range of 300–400K and found that the barrier height (~1.25 eV) and the ideality factor (~1.25) are nearly temperature independent. Similarly, Wang et al. [94] employed capacitance-voltage measurement technique to determine the Schottky barrier height to be 0.94 eV at room temperature. This section presents the study of the temperature-dependent electrical transport properties of InN/GaN heterostructures and observed the temperature-dependent barrier height and the ideality factor. Figure 9(a) shows the room temperature J-V(current density-voltage) characteristics for the junction. The junction between InN and GaN exhibits a rectifying behavior which suggests an existence of Schottky barrier height at the junction [95, 96].

To investigate further, we have studied the temperature-dependent J-V (J-V-T) characteristics ranging from 200 to 500 K and is given in Figure 9(b) [95]. In the present case, the GaN film is highly doped with silicon (ND ~ 1.4 1018 cm<sup>3</sup> ) resulting in a lower barrier height at the InN/ GaN interface, which is due to the tunneling of charge carriers through the interface. This kind of behavior reveals that the current transport is primarily dominated by thermionic field emission (TFE) mechanism, where the carriers tunnel from GaN to InN. According to the transport theory, TFE dominates only when E00/kT ≈ 1, where E00 is the characteristic tunneling energy which determines the nature of conduction mechanism. When E00/kT ≈ 1, the Boltzmann distribution tail of thermionic emission drops off by a factor of exp(1), which is much faster than the decrease rate of the tunneling probability. On the other hand, thermionic emission (TE) is predominant when E00/kT << 1 because the tunneling probability drops off faster than TE [97]. In the present case, the value of the barrier height and the ideality factor (η)

and asymmetric XRCs. The contrast at the interface in the TEM image in Figure 8(d) clearly indicates the presence of initial low-contrast Ga-rich layers which possibly attribute to the ~425 nm peak in PL spectra. The thickness was estimated to be ~100 nm from the TEM image. The other dark regions are formed as a result from the Ga ion beam damage during sample thinning.

Figure 8. (a) HRXRD 2θ-ω scan of InGaN on bare Si (111), (b) X-ray rocking curve of (0002) and (1011) reflections, (c) photoluminescence spectra of InGaN/Si, and (d) cross-sectional TEM of InGaN/Si heterostructure. Reprinted with per-

The semiconductor heterostructure exhibits the Schottky barrier at the interface due to the formation of conduction band offset because of their different bandgap values. The concepts on the band offset are directly transferable to the Schottky barrier height problems. Hence, it is

3. Transport properties of III-nitride heterostructure-based Schottky

3.1. Polar and nonpolar InN/GaN heterostructures.

junctions

mission from [88].

186 Epitaxy

Figure 9. (a) The room temperature J–V characteristics of InN/GaN Schottky junction and (b) the forward J–V characteristics with TFE fitting as a function of measurement temperature. Reprinted with permission from [95, 96].

were calculated by fitting a line in the linear region of the forward J-V curves using the TFE equation and are shown in Figure 9(b). From the abovementioned analysis, the barrier height and the ideality factor are found to be temperature dependent. Thus, our results indicate the presence of inhomogeneity at the interface due to the presence of various types of defects, which is why we observe the temperature-dependent behavior of barrier height [98, 99]. From the fitting the value of E00/kT is observed to be nearly one, suggesting that the TFE conduction mechanism would be considered to be a more realistic model for the analysis of the electronic transport in polar InN/GaN heterostructure.

characteristics of the InN/GaN heterostructure were measured at different temperatures. The values of the barrier height and ideality factor were estimated by fitting the linear region of the forward J-V curves using thermionic emission model. It was found that the ideality factor and the barrier height values range from 1.65 and 0.83 eV (500 K) to 4.1 and 0.4 eV (150 K), respectively, as shown in Figure 10(b). The variation of barrier height and the ideality factor with the measurement temperature indicates the presence of inhomogeneous nature of non-

The inhomogeneous nature of barrier height at the InN/GaN interface could arise due to the presence of various types of defects at the interface [98, 99]. The inhomogeneous nature at the interface was explained by using Richardson plot of saturation current. In the Richardson plot, we could identify two separate temperature ranges, i.e., 150–300 and 350–500 K from the slopes, and the values of Richardson's constants (A\*) were found to be much lower than the

the InN/GaN interface can be explained by considering the Gaussian distribution of barrier

deviation of barrier height, respectively. Considering Gaussian distribution of barrier height,

<sup>φ</sup><sup>b</sup> <sup>¼</sup> <sup>φ</sup>b<sup>0</sup> � <sup>q</sup>σ<sup>s</sup>

Here, φb<sup>0</sup> is the zero bias mean barrier height. Considering the barrier height inhomogeneities,

Figure 10(c) shows the modified Richardson plot. In the first region (300–500 K), the values of

constant value in the temperature range of 350–500 K is very close to the theoretical value of 24 A/cm<sup>2</sup> K<sup>2</sup> for n-type GaN. This indicates that at higher temperatures (350–500 K) the current transport is dominated by thermionic emission mechanism. The values of φ<sup>b</sup> and A\* for polar

The reduced barrier height in nonpolar a-plane InN/GaN heterojunction when compared to the polar c-plane InN/GaN can be attributed to the absence of polarization field at the interface. The value of A\* is largely deviated in the temperature range of 150–300 K indicating the reduced dominance of thermionic emission. Figure 10(d) shows E<sup>0</sup> = ηkT versus kT plot for the nonpolar a-InN/GaN Schottky diode. The value, E0, seems to be independent of temperature at

2

� <sup>q</sup><sup>2</sup>σ<sup>s</sup>

c-plane InN/GaN heterojunction were found to be 1.6 eV and 25.8 A/cm<sup>2</sup> K<sup>2</sup>

2k 2 T2 � �

<sup>2</sup><sup>π</sup> <sup>p</sup> exp � <sup>φ</sup><sup>b</sup> � <sup>φ</sup><sup>b</sup>

<sup>2</sup><sup>π</sup> <sup>p</sup> is the normalization constant and <sup>φ</sup><sup>b</sup> and <sup>σ</sup><sup>s</sup> are the mean and standard

2

<sup>¼</sup> ln AA� ð Þ� <sup>q</sup>φb<sup>0</sup>

K�<sup>2</sup> for n-GaN. This type of barrier height inhomogeneity at

Heterostructures of III-Nitride Semiconductors for Optical and Electronic Applications

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<sup>2</sup>kT (3)

kT (4)

, respectively [102].

, respectively. The calculated Richardson

(2)

189

� �<sup>2</sup> 2σ<sup>s</sup> 2 " #

polar InN/GaN interface.

theoretical value of 24 Acm�<sup>2</sup>

ffiffiffiffiffiffi

where 1=σ<sup>s</sup>

heights at the interface [100, 101] and can be written as

the effective barrier height, φb, given by the expression

the conventional Richardson plot is modified as follows:

<sup>φ</sup><sup>b</sup> and A\* were found to be 1.15 eV and 19.5 A/cm<sup>2</sup> <sup>K</sup><sup>2</sup>

ln Is T2 � �

P φ<sup>b</sup> � � <sup>¼</sup> <sup>1</sup> σs ffiffiffiffiffiffi

An optoelectronic device based on nonpolar III-nitride heterostructure has been an important subject due to its potential improvement on the efficiency. However, the transport behavior of nonpolar a-plane InN/GaN heterostructure interfaces is limited. In this section, we will discuss our results on the transport properties of nonpolar a-plane InN/GaN heterostructure. The inset of Figure 10(a) shows the schematic diagram of InN/GaN heterostructure Schottky junction. Figure 10(a) shows the room temperature I-V characteristic of nonpolar a-plane InN/GaN heterostructure Schottky junction [43]. The rectifying behavior of the I-V curve indicates the existence of Schottky barrier at the nonpolar InN/GaN interface. The forward bias J-V

Figure 10. (a) Room temperature I–V characteristics of the nonpolar a-plane InN/GaN heterostructures, (b) temperaturedependent barrier height and the ideality factor, (c) modified Richardson plot of ln(Is/T<sup>2</sup> )-q<sup>2</sup> σ2 /2k2 T2 versus q/kT, and (d) plot of ηkT as a function of kT. Reprinted with permission from [43, 96].

characteristics of the InN/GaN heterostructure were measured at different temperatures. The values of the barrier height and ideality factor were estimated by fitting the linear region of the forward J-V curves using thermionic emission model. It was found that the ideality factor and the barrier height values range from 1.65 and 0.83 eV (500 K) to 4.1 and 0.4 eV (150 K), respectively, as shown in Figure 10(b). The variation of barrier height and the ideality factor with the measurement temperature indicates the presence of inhomogeneous nature of nonpolar InN/GaN interface.

were calculated by fitting a line in the linear region of the forward J-V curves using the TFE equation and are shown in Figure 9(b). From the abovementioned analysis, the barrier height and the ideality factor are found to be temperature dependent. Thus, our results indicate the presence of inhomogeneity at the interface due to the presence of various types of defects, which is why we observe the temperature-dependent behavior of barrier height [98, 99]. From the fitting the value of E00/kT is observed to be nearly one, suggesting that the TFE conduction mechanism would be considered to be a more realistic model for the analysis of the electronic

An optoelectronic device based on nonpolar III-nitride heterostructure has been an important subject due to its potential improvement on the efficiency. However, the transport behavior of nonpolar a-plane InN/GaN heterostructure interfaces is limited. In this section, we will discuss our results on the transport properties of nonpolar a-plane InN/GaN heterostructure. The inset of Figure 10(a) shows the schematic diagram of InN/GaN heterostructure Schottky junction. Figure 10(a) shows the room temperature I-V characteristic of nonpolar a-plane InN/GaN heterostructure Schottky junction [43]. The rectifying behavior of the I-V curve indicates the existence of Schottky barrier at the nonpolar InN/GaN interface. The forward bias J-V

Figure 10. (a) Room temperature I–V characteristics of the nonpolar a-plane InN/GaN heterostructures, (b) temperature-

)-q<sup>2</sup> σ2 /2k2

T2 versus q/kT, and (d)

dependent barrier height and the ideality factor, (c) modified Richardson plot of ln(Is/T<sup>2</sup>

plot of ηkT as a function of kT. Reprinted with permission from [43, 96].

transport in polar InN/GaN heterostructure.

188 Epitaxy

The inhomogeneous nature of barrier height at the InN/GaN interface could arise due to the presence of various types of defects at the interface [98, 99]. The inhomogeneous nature at the interface was explained by using Richardson plot of saturation current. In the Richardson plot, we could identify two separate temperature ranges, i.e., 150–300 and 350–500 K from the slopes, and the values of Richardson's constants (A\*) were found to be much lower than the theoretical value of 24 Acm�<sup>2</sup> K�<sup>2</sup> for n-GaN. This type of barrier height inhomogeneity at the InN/GaN interface can be explained by considering the Gaussian distribution of barrier heights at the interface [100, 101] and can be written as

$$P(\boldsymbol{\varphi}\_b) = \frac{1}{\sigma\_s \sqrt{2\pi}} \exp\left[-\frac{\left(\boldsymbol{\varphi}\_b - \overline{\boldsymbol{\varphi}\_b}\right)^2}{2\sigma\_s^2}\right] \tag{2}$$

where 1=σ<sup>s</sup> ffiffiffiffiffiffi <sup>2</sup><sup>π</sup> <sup>p</sup> is the normalization constant and <sup>φ</sup><sup>b</sup> and <sup>σ</sup><sup>s</sup> are the mean and standard deviation of barrier height, respectively. Considering Gaussian distribution of barrier height, the effective barrier height, φb, given by the expression

$$
\overline{\rho}\_b = \overline{\rho\_{b0}} - \frac{q \sigma\_s^2}{2kT} \tag{3}
$$

Here, φb<sup>0</sup> is the zero bias mean barrier height. Considering the barrier height inhomogeneities, the conventional Richardson plot is modified as follows:

$$\ln\left(\frac{I\_s}{T^2}\right) - \left(\frac{q^2\sigma\_s^2}{2k^2T^2}\right) = \ln(AA^\*) - \frac{q\overline{q\rho\_{b0}}}{kT} \tag{4}$$

Figure 10(c) shows the modified Richardson plot. In the first region (300–500 K), the values of <sup>φ</sup><sup>b</sup> and A\* were found to be 1.15 eV and 19.5 A/cm<sup>2</sup> <sup>K</sup><sup>2</sup> , respectively. The calculated Richardson constant value in the temperature range of 350–500 K is very close to the theoretical value of 24 A/cm<sup>2</sup> K<sup>2</sup> for n-type GaN. This indicates that at higher temperatures (350–500 K) the current transport is dominated by thermionic emission mechanism. The values of φ<sup>b</sup> and A\* for polar c-plane InN/GaN heterojunction were found to be 1.6 eV and 25.8 A/cm<sup>2</sup> K<sup>2</sup> , respectively [102]. The reduced barrier height in nonpolar a-plane InN/GaN heterojunction when compared to the polar c-plane InN/GaN can be attributed to the absence of polarization field at the interface. The value of A\* is largely deviated in the temperature range of 150–300 K indicating the reduced dominance of thermionic emission. Figure 10(d) shows E<sup>0</sup> = ηkT versus kT plot for the nonpolar a-InN/GaN Schottky diode. The value, E0, seems to be independent of temperature at

temperature) and under ultraviolet radiation, which are consistent with the n-n isotype heterojunctions of other materials as reported by others [81, 82, 105]. I-V characteristics were obtained on Al/InGaN/Si(111)/Al, Al/InGaN/Al, and Al/Si (111)/Al. The behavior of Al/InGaN/ Al and Al/Si (111)/Al was ohmic and that of the Al/InGaN/Si (111)/Al junction was rectifying, thus confirming that the rectifying characteristic is primarily arising from the n-InGaN/n-Si isotype heterojunction. Although it seems to be like a leaky rectifying behavior, subsequent low-temperature current-voltage measurements were carried out to further confirm the rectifying behavior. Hall measurements were carried out, and negligible changes were observed at low temperatures in conductivity, mobility, carrier concentration, etc. The room temperature conductivity, mobility, and bulk carrier concentration of the InGaN layer were found to be

Heterostructures of III-Nitride Semiconductors for Optical and Electronic Applications

(~1019–1021) is a well-observed characteristic of the undoped InGaN and InN due to the nitrogen vacancies in the bulk and along the edge dislocations at the interface and In vacancy/N antisite complexes [106] which explains the source of doping in our case. A detailed

Although the device is not perfectly rectifying at room temperature, the interesting characteristics observed were in the region of zero bias. An abrupt increase in photocurrent in the presence of UV radiation was observed at zero bias than either reverse or positive biases and is shown in the inset of Figure 12(a). Ager et al. [83] have shown the operation of similar devices in photovoltaic mode. However, the InGaN films grown on their samples were much thicker compared to the present work and also used several buffer layers. The photocurrent response and stability were studied from the on-off cycles of a UV lamp at zero bias and different voltages and are shown in Figure 12(b). The order of increase in the photocurrent magnitude is higher in case of zero bias (>1.5) than that of the positive and reverse bias. Figure 12(c) and (d) shows the growth and decay responses

I tðÞ¼ Idark <sup>þ</sup> <sup>A</sup> <sup>1</sup> � exp �ð Þ <sup>t</sup> � <sup>t</sup><sup>0</sup>

I tðÞ¼ Idark <sup>þ</sup> <sup>A</sup> exp �ð Þ <sup>t</sup> � <sup>t</sup><sup>0</sup>

for growth and decay, respectively, where Idark is dark current; A is scaling constant; t<sup>0</sup> is the time when UV lamp was switched on or off for growth or decay, respectively; and τ<sup>g</sup> and τ<sup>d</sup> are growth and decay times. The response times obtained from the above equations are 20 and 33 ms for growth and decay. The responsivity (Rλ) of this photodetector is calculated from, Rλ=Iλ/PλS, where I<sup>λ</sup> is photocurrent, P<sup>λ</sup> is the incident power of UV lamp of wavelength, and S is the area

EQE = hcRλ/eλ, [108]. The values obtained for spectral responsivity and external quantum efficiency are 0.0942 A/W and 32.4%, respectively, which are better than the values reported in the literature [85] for such devices. The responsivities and external quantum efficiencies in case of different biases were found to be better than zero bias (Table 3). From Figure 12(b), it can be seen that the photocurrents obtained at different biases do not overlap, i.e., the photocurrent

τg

τg

(5)

(6)

. The external quantum efficiency (EQE) is given as

. The high background electron concentration

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191

~201 (Ω cm)�<sup>1</sup>

, ~17 cm2

which can be described as [107]

of illuminated junction which is 0.09 cm<sup>2</sup>

/Vs, and ~8 � 1019 cm�<sup>3</sup>

study on the current-transport mechanisms can be found elsewhere [88].

Figure 11. (a) Room temperature J-V characteristics of the InGaN/GaN heterostructure, (b) temperature-dependent J-V characteristics of the InGaN/GaN heterostructure, and (c) variation of the barrier height and the ideality factor with measurement temperature. Reprinted with permission from [104].

low temperatures, thus indicating the dominance of field emission in the range of 150–500 K [103]. It can be understood that the carriers lack the required energy to surmount the barrier at the low temperature and thus tunnel through the defect states at the interface.

#### 3.2. Polar InGaN/GaN heterostructures

In this section, we have grown InGaN/GaN heterostructures using plasma-assisted molecular beam epitaxy and studied the temperature-dependent electrical transport characteristics. Figure 11(a) shows the room temperature J-V characteristics of the InGaN/GaN heterostructure in both linear and semilog scale [104]. The schematic diagram of the device structure has been shown in the inset of Figure 11(a). The rectifying nature of J-V characteristic of InGaN/GaN heterostructure suggested the presence of a Schottky barrier at the interface. The temperaturedependent J-V characteristic of the heterostructure is shown in Figure 11(b). It is observed that, as the measurement temperature increases, the forward bias current increases, which indicate that the current transport across the Schottky junction is governed by the thermionic emission mechanism. Then, the values of barrier height and ideality factor were calculated by using thermionic emission model. Figure 11(c) shows the variation of barrier height and ideality factor with measurement temperature [104]. It is observed that there is a temperature-dependent variation of both barrier height and ideality factor. The temperature dependence of φ<sup>b</sup> indicates that the barrier height is inhomogeneous in nature, which may be due to various types of defects present at the InGaN/GaN interface. Moreover, the observed ideality factor greater than unity indicates a nonideal nature of the heterojunction, which is attributed to the presence of interface defect states.

#### 4. InGaN/Si heterostructure-based UV and IR photodetectors

Room temperature I-V measurements were performed on InGaN/Si (111) heterojunctions with Si biased positively and were shown in Figure 3(a). From Figure 3(a), it is observed that the device is showing rectifying characteristics, both in the dark (room temperature as well as low temperature) and under ultraviolet radiation, which are consistent with the n-n isotype heterojunctions of other materials as reported by others [81, 82, 105]. I-V characteristics were obtained on Al/InGaN/Si(111)/Al, Al/InGaN/Al, and Al/Si (111)/Al. The behavior of Al/InGaN/ Al and Al/Si (111)/Al was ohmic and that of the Al/InGaN/Si (111)/Al junction was rectifying, thus confirming that the rectifying characteristic is primarily arising from the n-InGaN/n-Si isotype heterojunction. Although it seems to be like a leaky rectifying behavior, subsequent low-temperature current-voltage measurements were carried out to further confirm the rectifying behavior. Hall measurements were carried out, and negligible changes were observed at low temperatures in conductivity, mobility, carrier concentration, etc. The room temperature conductivity, mobility, and bulk carrier concentration of the InGaN layer were found to be ~201 (Ω cm)�<sup>1</sup> , ~17 cm2 /Vs, and ~8 � 1019 cm�<sup>3</sup> . The high background electron concentration (~1019–1021) is a well-observed characteristic of the undoped InGaN and InN due to the nitrogen vacancies in the bulk and along the edge dislocations at the interface and In vacancy/N antisite complexes [106] which explains the source of doping in our case. A detailed study on the current-transport mechanisms can be found elsewhere [88].

low temperatures, thus indicating the dominance of field emission in the range of 150–500 K [103]. It can be understood that the carriers lack the required energy to surmount the barrier at

Figure 11. (a) Room temperature J-V characteristics of the InGaN/GaN heterostructure, (b) temperature-dependent J-V characteristics of the InGaN/GaN heterostructure, and (c) variation of the barrier height and the ideality factor with

In this section, we have grown InGaN/GaN heterostructures using plasma-assisted molecular beam epitaxy and studied the temperature-dependent electrical transport characteristics. Figure 11(a) shows the room temperature J-V characteristics of the InGaN/GaN heterostructure in both linear and semilog scale [104]. The schematic diagram of the device structure has been shown in the inset of Figure 11(a). The rectifying nature of J-V characteristic of InGaN/GaN heterostructure suggested the presence of a Schottky barrier at the interface. The temperaturedependent J-V characteristic of the heterostructure is shown in Figure 11(b). It is observed that, as the measurement temperature increases, the forward bias current increases, which indicate that the current transport across the Schottky junction is governed by the thermionic emission mechanism. Then, the values of barrier height and ideality factor were calculated by using thermionic emission model. Figure 11(c) shows the variation of barrier height and ideality factor with measurement temperature [104]. It is observed that there is a temperature-dependent variation of both barrier height and ideality factor. The temperature dependence of φ<sup>b</sup> indicates that the barrier height is inhomogeneous in nature, which may be due to various types of defects present at the InGaN/GaN interface. Moreover, the observed ideality factor greater than unity indicates a nonideal nature of the heterojunction, which is attributed to the presence of interface

the low temperature and thus tunnel through the defect states at the interface.

4. InGaN/Si heterostructure-based UV and IR photodetectors

Room temperature I-V measurements were performed on InGaN/Si (111) heterojunctions with Si biased positively and were shown in Figure 3(a). From Figure 3(a), it is observed that the device is showing rectifying characteristics, both in the dark (room temperature as well as low

3.2. Polar InGaN/GaN heterostructures

measurement temperature. Reprinted with permission from [104].

defect states.

190 Epitaxy

Although the device is not perfectly rectifying at room temperature, the interesting characteristics observed were in the region of zero bias. An abrupt increase in photocurrent in the presence of UV radiation was observed at zero bias than either reverse or positive biases and is shown in the inset of Figure 12(a). Ager et al. [83] have shown the operation of similar devices in photovoltaic mode. However, the InGaN films grown on their samples were much thicker compared to the present work and also used several buffer layers. The photocurrent response and stability were studied from the on-off cycles of a UV lamp at zero bias and different voltages and are shown in Figure 12(b). The order of increase in the photocurrent magnitude is higher in case of zero bias (>1.5) than that of the positive and reverse bias. Figure 12(c) and (d) shows the growth and decay responses which can be described as [107]

$$I(t) = I\_{dark} + A \left[ 1 - \exp\left\{ \frac{-(t - t\_0)}{\tau\_\mathcal{g}} \right\} \right] \tag{5}$$

$$I(t) = I\_{dark} + A \left[ \exp\left\{ \frac{-(t - t\_0)}{\tau\_\mathcal{g}} \right\} \right] \tag{6}$$

for growth and decay, respectively, where Idark is dark current; A is scaling constant; t<sup>0</sup> is the time when UV lamp was switched on or off for growth or decay, respectively; and τ<sup>g</sup> and τ<sup>d</sup> are growth and decay times. The response times obtained from the above equations are 20 and 33 ms for growth and decay. The responsivity (Rλ) of this photodetector is calculated from, Rλ=Iλ/PλS, where I<sup>λ</sup> is photocurrent, P<sup>λ</sup> is the incident power of UV lamp of wavelength, and S is the area of illuminated junction which is 0.09 cm<sup>2</sup> . The external quantum efficiency (EQE) is given as EQE = hcRλ/eλ, [108]. The values obtained for spectral responsivity and external quantum efficiency are 0.0942 A/W and 32.4%, respectively, which are better than the values reported in the literature [85] for such devices. The responsivities and external quantum efficiencies in case of different biases were found to be better than zero bias (Table 3). From Figure 12(b), it can be seen that the photocurrents obtained at different biases do not overlap, i.e., the photocurrent

height and resulting in abrupt increase of electron flow. However, in case of a bias (positive or negative as explained above), the electrons at the interface tunnel through the depletion region leaving behind holes which are eventually refilled by electrons from the other side. The forward and negative bias characteristics from the current-voltage plots are in good agreement

Heterostructures of III-Nitride Semiconductors for Optical and Electronic Applications

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193

In summary, we presented the studies on the growth, characterizations, and transport properties of III-nitride-based heterostructures. Here, discussion has been carried out on the growth of polar InN/GaN, nonpolar InN/GaN, polar InGaN/GaN, nonpolar InGaN/GaN, and InGaN/Si heterostructures by using MBE system followed by their characterizations. Moreover, we have presented the electrical transport properties across the heterostructures interface. In addition, UV and IR photodetection studies on InGaN/Si heterostructures have

[1] Yoshida S, Misawa S, Gonda S. Improvements on the electrical and luminescent properties of reactive molecular beam epitaxially grown GaN films by using AlN‐coated

[2] Akasaki I, Amano H, Koide Y, Hiramatsu H, Sawak N. Effects of ain buffer layer on crystallographic structure and on electrical and optical properties of GaN and Ga<sup>1</sup><sup>x</sup> AlxN (0 < x ≦ 0.4) films grown on sapphire substrate by MOVPE. Journal of Crystal

[3] Nakamura S. GaN growth using GaN buffer layer. Japanese Journal of Applied Physics.

[4] Nakamura S, Senoh M, Mukai T. P-GaN/N-InGaN/N-GaN double-heterostructure blue-

light-emitting diodes. Japanese Journal of Applied Physics. 1993;32:L8-L11

, Shruti Mukundan<sup>1</sup> and Saluru Baba Krupanidhi<sup>1</sup>

\*

with the proposed mechanism.

5. Conclusions

been discussed.

Author details

References

Basanta Roul1,2, Greeshma Chandan1

Growth. 1989;98:209

1991;30:L1705

\*Address all correspondence to: sbkenator@gmail.com

1 Materials Research Centre, Indian Institute of Science, Bangalore, India

2 Central Research Laboratory, Bharat Electronics, Bangalore, India

sapphire substrates. Applied Physics Letters. 1983;42:427

Figure 12. (a) Room temperature current-voltage characteristics of n-InGaN/n-Si heterojunction in dark at room temperature and at 133 K and under UV radiation at room temperature, (b) photoresponse at different working voltages, (c) time response of photocurrent growth, and (d) decay from the fitting of experimental values. Reprinted with permission from [88].


Table 3. Bias-dependent responsivities and external quantum efficiencies.

range in each case is distinct from that of others, suggesting that such devices can be used for switching purposes and logical operations as well.

The mechanism for this type of behavior in n-n isotype heterojunctions can be explained with the help of a model proposed by Yawata et al. [77]. When the electron and hole pairs are generated, the electrons are swept away from the junctions due to the built-in electric field, whereas the holes are trapped in the notch. The holes being positively charged neutralize the electrons at interface and eventually lower the barrier height. For an intense illumination, the concentration of holes trapped at the notch is increased, thus drastically lowering the barrier height and resulting in abrupt increase of electron flow. However, in case of a bias (positive or negative as explained above), the electrons at the interface tunnel through the depletion region leaving behind holes which are eventually refilled by electrons from the other side. The forward and negative bias characteristics from the current-voltage plots are in good agreement with the proposed mechanism.

## 5. Conclusions

In summary, we presented the studies on the growth, characterizations, and transport properties of III-nitride-based heterostructures. Here, discussion has been carried out on the growth of polar InN/GaN, nonpolar InN/GaN, polar InGaN/GaN, nonpolar InGaN/GaN, and InGaN/Si heterostructures by using MBE system followed by their characterizations. Moreover, we have presented the electrical transport properties across the heterostructures interface. In addition, UV and IR photodetection studies on InGaN/Si heterostructures have been discussed.

## Author details

Basanta Roul1,2, Greeshma Chandan1 , Shruti Mukundan<sup>1</sup> and Saluru Baba Krupanidhi<sup>1</sup> \*

\*Address all correspondence to: sbkenator@gmail.com


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range in each case is distinct from that of others, suggesting that such devices can be used for

Figure 12. (a) Room temperature current-voltage characteristics of n-InGaN/n-Si heterojunction in dark at room temperature and at 133 K and under UV radiation at room temperature, (b) photoresponse at different working voltages, (c) time response of photocurrent growth, and (d) decay from the fitting of experimental values. Reprinted with permission from [88].

Voltage (V) Responsivity (A/W) EQE (%) 0.0942 32.4 0.6217 213.8 1.7097 588.1 0.5746 197.6 1.2575 432.5

The mechanism for this type of behavior in n-n isotype heterojunctions can be explained with the help of a model proposed by Yawata et al. [77]. When the electron and hole pairs are generated, the electrons are swept away from the junctions due to the built-in electric field, whereas the holes are trapped in the notch. The holes being positively charged neutralize the electrons at interface and eventually lower the barrier height. For an intense illumination, the concentration of holes trapped at the notch is increased, thus drastically lowering the barrier

switching purposes and logical operations as well.

192 Epitaxy

Table 3. Bias-dependent responsivities and external quantum efficiencies.


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**Chapter 9**

**Provisional chapter**

Ga1−xAs/InP (0.53 < x < 1)

**Epitaxy and Device Properties of InGaAs**

**Epitaxy and Device Properties of InGaAs** 

Xing-you Chen, Yi Gu and Yong-gang Zhang

Xing-you Chen, Yi Gu and Yong-gang Zhang

Additional information is available at the end of the chapter

structures on lattice-mismatched material systems.

mature with cutoff wavelength at 1.7 μm. Wavelength-extended In<sup>x</sup>

force microscopy, photoluminescence

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/intechopen.70259

**Abstract**

**1. Introduction**

**Photodetectors with Relatively High Lattice Mismatch**

In this chapter, our works on the developments of wavelength-extended InGaAs photodetectors with cutoff wavelength >1.7 μm are reviewed. Various InGaAs/InAlAs p-i-n heterojunction structures have been grown on InP and GaAs substrates by gas source molecular beam epitaxy, some details on the InGaAs photodetector structures and the techniques of metamorphic buffer layer such as linearly, step, and one-step continuously InAlAs graded buffer, and dislocation restraint methods of compositional overshoot and digital alloy are introduced. The material characteristics and device properties were evaluated by atomic force microscopy, high-resolution X-ray diffraction and reciprocal space mapping, cross-sectional transmission electron microscopy, and current-voltage measurements, etc. The results provide clues to the development of metamorphic device

**Keywords:** InGaAs, photodetector, metamorphic, lattice mismatch, X-ray diffraction, atomic

InGaAs photodetectors (PDs) and focal plane arrays (FPAs) are attracting particular interests as they can be tailored to cover one of the atmospheric windows of 1–3 μm in shortwave infrared band. The ln0.53Ga0.47As PDs grown lattice-matched to InP are commercially

PDs with cutoff wavelength more than 1.7 μm have been extensively investigated over the past decades due to their important applications in spatial remote sensing, earth observation, environmental monitoring, etc. [1, 2]. However, for application in longer wavelength region, the indium content should be increased, which increases the lattice mismatch

**Photodetectors with Relatively High Lattice Mismatch**

DOI: 10.5772/intechopen.70259

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution,

© 2018 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

and reproduction in any medium, provided the original work is properly cited.

**Provisional chapter**

## **Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch**

DOI: 10.5772/intechopen.70259

Xing-you Chen, Yi Gu and Yong-gang Zhang Xing-you Chen, Yi Gu and Yong-gang Zhang

Additional information is available at the end of the chapter Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/intechopen.70259

#### **Abstract**

In this chapter, our works on the developments of wavelength-extended InGaAs photodetectors with cutoff wavelength >1.7 μm are reviewed. Various InGaAs/InAlAs p-i-n heterojunction structures have been grown on InP and GaAs substrates by gas source molecular beam epitaxy, some details on the InGaAs photodetector structures and the techniques of metamorphic buffer layer such as linearly, step, and one-step continuously InAlAs graded buffer, and dislocation restraint methods of compositional overshoot and digital alloy are introduced. The material characteristics and device properties were evaluated by atomic force microscopy, high-resolution X-ray diffraction and reciprocal space mapping, cross-sectional transmission electron microscopy, and current-voltage measurements, etc. The results provide clues to the development of metamorphic device structures on lattice-mismatched material systems.

**Keywords:** InGaAs, photodetector, metamorphic, lattice mismatch, X-ray diffraction, atomic force microscopy, photoluminescence

## **1. Introduction**

InGaAs photodetectors (PDs) and focal plane arrays (FPAs) are attracting particular interests as they can be tailored to cover one of the atmospheric windows of 1–3 μm in shortwave infrared band. The ln0.53Ga0.47As PDs grown lattice-matched to InP are commercially mature with cutoff wavelength at 1.7 μm. Wavelength-extended In<sup>x</sup> Ga1−xAs/InP (0.53 < x < 1) PDs with cutoff wavelength more than 1.7 μm have been extensively investigated over the past decades due to their important applications in spatial remote sensing, earth observation, environmental monitoring, etc. [1, 2]. However, for application in longer wavelength region, the indium content should be increased, which increases the lattice mismatch

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2018 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

between high indium content InGaAs and InP substrate. For instance, the In0.83Ga0.17As PD with cutoff wavelength of 2.5 μm possesses a lattice mismatch of about +2.1% with respect to InP. In that case, the dark current increases several orders of magnitude, which severely hinders the development of wavelength-extended InGaAs PDs in the infrared range. To suppress the dark current of InGaAs PD, it is a prerequisite to optimize the InGaAs material with high indium (In) content. Therefore, high In content In<sup>x</sup> Ga1−xAs has been grown on InP, GaAs, and Si substrate by using techniques such as molecular beam epitaxy (MBE) and metal organic chemical vapor deposition (MOCVD). During the process of material growth, dislocation restriction techniques such as beryllium (Be) doping [3], dilute nitride [4], strained or strain-compensated superlattice (SL), and strain-driven quantum dots (QDs) were adopted to reduce the threading dislocation density (TDD) in the metamorphic buffer layer (MBL). Several sets of results with different x values of the In<sup>x</sup> Ga1−xAs layers are listed in **Table 1**. Our own results are also included. It can be seen that the TDD could be quite different for InGaAs with various In contents due to different lattice mismatches with different substrates.

been optimized to acquire high-performance InGaAs PDs with relatively high-lattice mismatch. Additionally, some strategies were used to reduce the residual strain and decrease the TDD, such as composition overshoot and the insertion of digital alloy (DA) intermediate layer in the MBL. All the material growth in this chapter was performed in a VG Semicon V80H gas source molecular beam epitaxy (GSMBE) system. The best background chamber pressure achieved in this system was about 3 × 10−11 Torr, and the pressure during growth is typically in the 10−5 Torr range. The elemental In, gallium (Ga), and aluminum (Al) sources were used as group III sources, and the elemental silicon (Si) and Be were used as n-type and p-type dopant sources, respectively. Their fluxes were controlled by changing the cell temperatures. Arsine and phosphine cracking cells were used as group V sources, and their fluxes were controlled by adjusting the pressure. The cracking temperature was approximately 1000°C measured by using a thermocouple. The growth rates were all adjusted to about 1 μm/h for InP, InAlAs,

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch

http://dx.doi.org/10.5772/intechopen.70259

205

**2. Optimize the growth of wavelength extended InGaAs PDs on InP**

perature should be optimized for each layer. Generally, a high *T*<sup>g</sup>

increase the dark current of PDs. In this work, *T*<sup>g</sup>

= 3 × 1018 cm−3) InAlAs buffer layer and a 2-μm-thick n−

and InGaAs absorption layer, respectively [17].

low, whereas the InAlAs MBL with the *T*<sup>g</sup>

**2.1. The effects of growth temperature on the characteristics of InP-based In0.83Ga0.17As PDs**

high-device performance wavelength-extended InGaAs PD on InP substrate, the growth tem-

ration of background impurities into the epilayers during the growth process, whereas the

which will cause a poor surface and nonuniformity of compositions in the epilayer and then

Four In0.83Ga0.17As PD structures with continuously graded InAlAs buffers were grown on (1 0 0)-oriented Fe-doped InP epi-ready substrate. Each structure consisted of a 1.95-μm N+

The results show that the growth temperature of the absorption layer should not be very

talline quality than that grown at either 460 or 500°C. The sample number 3 with InAlAs MBL grown with temperature graded from 500 to 460°C, and InGaAs absorption layer grown at 500°C exhibits better crystalline quality. It shows smoother surface with lower root-mean-square (RMS) of atomic force microscopy (AFM), narrower full width at half maximum (FWHM) of X-ray diffraction (XRD), and stronger photoluminescence (PL) intensity than others, as shown in **Figure 2**. Therefore, this temperature strategy for growth of wavelength-extended InGaAs PDs on InP substrate will be always used hereafter except for

at stages A, B, C, and D of the four PD structures were shown in **Figure 1**.

will enhance the In segregation on the surface of InGaAs with a high In content [16],

) is one of the most important parameters of MBE. To acquire

(P = 7 × 1018 cm−3) 600-nm-thick In0.84Al0.16As cap. The In profiles and

can reduce the incorpo-

(N

was optimized for the InAlAs buffer layer

(n = 3 × 1016 cm−3) In0.84Ga0.16As absorp-

graded from 500 to 460°C exhibits better crys-

and InGaAs.

higher *T*<sup>g</sup>

*T*g

The growth temperature (*T*<sup>g</sup>

tion layer followed by a P<sup>+</sup>

special reminding.

In this chapter, the InGaAs/InAlAs p-i-n heterojunction structure was always used for the growth of wavelength-extended InGaAs PD with high indium content because of higher quantum efficiency than homostructure. Both the MBL schemes and experimental parameters have


**Table 1.** Parameters of Inx Ga1−xAs (0.53 ≤ x ≤ 1) as reported by various researchers. been optimized to acquire high-performance InGaAs PDs with relatively high-lattice mismatch. Additionally, some strategies were used to reduce the residual strain and decrease the TDD, such as composition overshoot and the insertion of digital alloy (DA) intermediate layer in the MBL.

between high indium content InGaAs and InP substrate. For instance, the In0.83Ga0.17As PD with cutoff wavelength of 2.5 μm possesses a lattice mismatch of about +2.1% with respect to InP. In that case, the dark current increases several orders of magnitude, which severely hinders the development of wavelength-extended InGaAs PDs in the infrared range. To suppress the dark current of InGaAs PD, it is a prerequisite to optimize the InGaAs mate-

on InP, GaAs, and Si substrate by using techniques such as molecular beam epitaxy (MBE) and metal organic chemical vapor deposition (MOCVD). During the process of material growth, dislocation restriction techniques such as beryllium (Be) doping [3], dilute nitride [4], strained or strain-compensated superlattice (SL), and strain-driven quantum dots (QDs) were adopted to reduce the threading dislocation density (TDD) in the metamorphic buf-

listed in **Table 1**. Our own results are also included. It can be seen that the TDD could be quite different for InGaAs with various In contents due to different lattice mismatches with

In this chapter, the InGaAs/InAlAs p-i-n heterojunction structure was always used for the growth of wavelength-extended InGaAs PD with high indium content because of higher quantum efficiency than homostructure. Both the MBL schemes and experimental parameters have

–10<sup>9</sup>

0.82 MOCVD InP (0 0 1) ~1011–1012 (XRD-FWHM) Zhao et al. [7]

0.85 SSMBE GaAs (0 0 1) Not mentioned Jurczak et al. [14]

0.53 SSMBE Si (1 1 1) Not mentioned Gao et al. [15]

Ga1−xAs (0.53 ≤ x ≤ 1) as reported by various researchers.

Ga1−xAs has been grown

Ga1−xAs layers are

**Ga1−xAs (cm−2) Reference**

(XTEM) Ji et al. [5]

(XTEM) Zhao et al. [6]

(XTEM) Present work

(XTEM) Lubyshev et al. [8]

(XTEM) Valtuefia et al. [9]

–1010 (plan-view TEM) Chang et al. [10]

(XTEM) Song et al. [13]

–1010 (XTEM) Present work

(XTEM) Chang et al. [11]

(EPD) Zimmermann et al. [12]

rial with high indium (In) content. Therefore, high In content In<sup>x</sup>

fer layer (MBL). Several sets of results with different x values of the In<sup>x</sup>

**X value Growth method Substrate TDD in Inx**

0.68 MOCVD InP (0 0 1) ~106

0.82 MOCVD InP (0 0 1) ~108

0.83 GSMBE InP (0 0 1) ≤107

0.53 SSMBE GaAs (0 0 1) ~106

0.6 SSMBE GaAs (0 0 1) ~108

0.75–1 SSMBE GaAs (0 0 1) ~10<sup>9</sup>

1 SSMBE GaAs (0 0 1) ~108

0.8 Not mentioned GaAs (0 0 1) ~105

0.63 SSMBE GaAs (0 0 1) ~108

0.83 GSMBE GaAs (0 0 1) ~10<sup>9</sup>

**Table 1.** Parameters of Inx

different substrates.

204 Epitaxy

All the material growth in this chapter was performed in a VG Semicon V80H gas source molecular beam epitaxy (GSMBE) system. The best background chamber pressure achieved in this system was about 3 × 10−11 Torr, and the pressure during growth is typically in the 10−5 Torr range. The elemental In, gallium (Ga), and aluminum (Al) sources were used as group III sources, and the elemental silicon (Si) and Be were used as n-type and p-type dopant sources, respectively. Their fluxes were controlled by changing the cell temperatures. Arsine and phosphine cracking cells were used as group V sources, and their fluxes were controlled by adjusting the pressure. The cracking temperature was approximately 1000°C measured by using a thermocouple. The growth rates were all adjusted to about 1 μm/h for InP, InAlAs, and InGaAs.

## **2. Optimize the growth of wavelength extended InGaAs PDs on InP**

## **2.1. The effects of growth temperature on the characteristics of InP-based In0.83Ga0.17As PDs**

The growth temperature (*T*<sup>g</sup> ) is one of the most important parameters of MBE. To acquire high-device performance wavelength-extended InGaAs PD on InP substrate, the growth temperature should be optimized for each layer. Generally, a high *T*<sup>g</sup> can reduce the incorporation of background impurities into the epilayers during the growth process, whereas the higher *T*<sup>g</sup> will enhance the In segregation on the surface of InGaAs with a high In content [16], which will cause a poor surface and nonuniformity of compositions in the epilayer and then increase the dark current of PDs. In this work, *T*<sup>g</sup> was optimized for the InAlAs buffer layer and InGaAs absorption layer, respectively [17].

Four In0.83Ga0.17As PD structures with continuously graded InAlAs buffers were grown on (1 0 0)-oriented Fe-doped InP epi-ready substrate. Each structure consisted of a 1.95-μm N+ (N = 3 × 1018 cm−3) InAlAs buffer layer and a 2-μm-thick n− (n = 3 × 1016 cm−3) In0.84Ga0.16As absorption layer followed by a P<sup>+</sup> (P = 7 × 1018 cm−3) 600-nm-thick In0.84Al0.16As cap. The In profiles and *T*g at stages A, B, C, and D of the four PD structures were shown in **Figure 1**.

The results show that the growth temperature of the absorption layer should not be very low, whereas the InAlAs MBL with the *T*<sup>g</sup> graded from 500 to 460°C exhibits better crystalline quality than that grown at either 460 or 500°C. The sample number 3 with InAlAs MBL grown with temperature graded from 500 to 460°C, and InGaAs absorption layer grown at 500°C exhibits better crystalline quality. It shows smoother surface with lower root-mean-square (RMS) of atomic force microscopy (AFM), narrower full width at half maximum (FWHM) of X-ray diffraction (XRD), and stronger photoluminescence (PL) intensity than others, as shown in **Figure 2**. Therefore, this temperature strategy for growth of wavelength-extended InGaAs PDs on InP substrate will be always used hereafter except for special reminding.


**Figure 1.** Growth temperatures at points A–D of the In0.84Ga0.16As PD structures. The inset shows the indium composition vs. the grown thickness of samples 1–4. Reprinted with permission from Elsevier.

#### **2.2. The optimization and comparison of PD structures on different metamorphic buffer layers**

To acquire high-quality wavelength-extended InGaAs materials with relatively high-lattice mismatch, an appropriate InAlAs MBL is needed. Generally, for the growth of latticemismatched material system, the MBLs should be thick enough so that the lattice parameter can finally relax to be a freestanding status, and the misfit dislocation can annihilate ultimately during the growth process of the MBL. However, it is impractical and high cost to grow an excessively thick MBL for device structure by MBE technique due to a limited growth rate of around 1 μm h−1. Therefore, proper schemes of thin MBL should be explored.

to 0.83 through the simultaneously linear increase of In source temperature and decrease of Al source temperature in a relatively small temperature region. This caused the In and Al beam fluxes to change with cell temperatures exponentially in opposite directions so that the composition grading profile in the buffer approximates a linear grade. The thickness of the InAlAs MBL was about 1.9 μm. After that, the growth temperature was increased to 490°C for

**Figure 2.** Summary of the measured data of samples with different temperature strategies. (a) HRXRD FWHM of substrate peak, (b) HRXRD FWHM of layer peak, (c) PL intensity of InGaAs absorption layer and (d) AFM RMS of

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch

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207

Al1−xAs buffer layers with In composition of 0.60, 0.68, 0.76, and 0.83 were grown in

sequence at 490, 480, 470, and 460°C, respectively. The In composition x was step graded from 0.52 to 0.83. Therefore, the thickness of the step-graded InAlAs MBL was also about 1.9 μm with each step layer of about 475 nm. After that, the growth temperature was kept at 490°C

To reduce the conduction band discontinuity and the dislocation density between the InGaAs absorption layer and InAlAs buffer and cap layers due to the small lattice mismatch between

Al1−xAs MBL started with a 200-nm InP buffer

In0.52Al0.48As layer grown at 500°C, and then

the following growth.

Inx

for the following growth.

four N<sup>+</sup>

The growth of sample S2 with a step-graded In<sup>x</sup>

samples 1–4. Reprinted with permission from Elsevier.

layer grown at 460°C, following a 0.1-μm-thick N+

In our previous work, the effect of various kinds of MBL schemes such as a thick uniform buffer, continuously graded buffer, and step-graded buffer on the structural characteristics and device performances of InP-based InGaAs PDs has been compared [18–20].

Two In0.83Ga0.17As PD wafers with continuously (sample S1) and step-graded (sample S2) InAlAs MBLs were grown on semi-insulated (S.I.) (1 0 0)-oriented InP epi-ready substrate. Each wafer consisted of a 1.9-μm N+ (N = 3 × 1018 cm−3) InAlAs MBL and a 1.5-μm-thick n− (n = 3 × 1016 cm−3) In0.83Ga0.17As absorption layer followed by a P<sup>+</sup> (P = 7 × 1018 cm−3) 530-nm-thick InAlAs cap layer. The typical schematic structures of the two PDs were shown in **Figure 3**.

The growth of sample S1 with a continuously graded In<sup>x</sup> Al1−xAs MBL started with a 200 nm InP buffer layer grown at 460°C, following a 0.1-μm-thick N+ In0.52Al0.48As layer grown at 500°C, the N+ continuously graded Inx Al1−xAs buffer layer was grown with growth temperature decreased linearly from 500 to 460°C, and the In composition x was graded from 0.52

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch http://dx.doi.org/10.5772/intechopen.70259 207

**Figure 2.** Summary of the measured data of samples with different temperature strategies. (a) HRXRD FWHM of substrate peak, (b) HRXRD FWHM of layer peak, (c) PL intensity of InGaAs absorption layer and (d) AFM RMS of samples 1–4. Reprinted with permission from Elsevier.

**2.2. The optimization and comparison of PD structures on different metamorphic buffer** 

**Figure 1.** Growth temperatures at points A–D of the In0.84Ga0.16As PD structures. The inset shows the indium composition

To acquire high-quality wavelength-extended InGaAs materials with relatively high-lattice mismatch, an appropriate InAlAs MBL is needed. Generally, for the growth of latticemismatched material system, the MBLs should be thick enough so that the lattice parameter can finally relax to be a freestanding status, and the misfit dislocation can annihilate ultimately during the growth process of the MBL. However, it is impractical and high cost to grow an excessively thick MBL for device structure by MBE technique due to a limited growth rate of around 1 μm h−1. Therefore, proper schemes of thin MBL should be explored. In our previous work, the effect of various kinds of MBL schemes such as a thick uniform buffer, continuously graded buffer, and step-graded buffer on the structural characteristics and

Two In0.83Ga0.17As PD wafers with continuously (sample S1) and step-graded (sample S2) InAlAs MBLs were grown on semi-insulated (S.I.) (1 0 0)-oriented InP epi-ready substrate.

InAlAs cap layer. The typical schematic structures of the two PDs were shown in **Figure 3**.

ture decreased linearly from 500 to 460°C, and the In composition x was graded from 0.52

(N = 3 × 1018 cm−3) InAlAs MBL and a 1.5-μm-thick n−

Al1−xAs buffer layer was grown with growth tempera-

(P = 7 × 1018 cm−3) 530-nm-thick

Al1−xAs MBL started with a 200-

In0.52Al0.48As layer grown at

(n

device performances of InP-based InGaAs PDs has been compared [18–20].

= 3 × 1016 cm−3) In0.83Ga0.17As absorption layer followed by a P<sup>+</sup>

vs. the grown thickness of samples 1–4. Reprinted with permission from Elsevier.

The growth of sample S1 with a continuously graded In<sup>x</sup>

continuously graded Inx

nm InP buffer layer grown at 460°C, following a 0.1-μm-thick N+

Each wafer consisted of a 1.9-μm N+

500°C, the N+

**layers**

206 Epitaxy

to 0.83 through the simultaneously linear increase of In source temperature and decrease of Al source temperature in a relatively small temperature region. This caused the In and Al beam fluxes to change with cell temperatures exponentially in opposite directions so that the composition grading profile in the buffer approximates a linear grade. The thickness of the InAlAs MBL was about 1.9 μm. After that, the growth temperature was increased to 490°C for the following growth.

The growth of sample S2 with a step-graded In<sup>x</sup> Al1−xAs MBL started with a 200-nm InP buffer layer grown at 460°C, following a 0.1-μm-thick N+ In0.52Al0.48As layer grown at 500°C, and then four N<sup>+</sup> Inx Al1−xAs buffer layers with In composition of 0.60, 0.68, 0.76, and 0.83 were grown in sequence at 490, 480, 470, and 460°C, respectively. The In composition x was step graded from 0.52 to 0.83. Therefore, the thickness of the step-graded InAlAs MBL was also about 1.9 μm with each step layer of about 475 nm. After that, the growth temperature was kept at 490°C for the following growth.

To reduce the conduction band discontinuity and the dislocation density between the InGaAs absorption layer and InAlAs buffer and cap layers due to the small lattice mismatch between

**Figure 3.** The schematic structure of In0.83Ga0.17As photodetectors.

InGaAs and InAlAs at the interfaces, digital-graded superlattices (DGSLs) were grown at the interfaces, respectively. First, DGSL1 consisted of nine groups of In0.83Ga0.17As/In0.83Al0.17As thin films which alternate with each other. The thickness of each group is about 8 nm. The thickness ratios of In0.83Ga0.17As/In0.83Al0.17As in each group are 1:9, 2:8, …, finally coming to …, 8:2, 9:1, respectively. Similarly, DGSL2 was comprised of nine groups of In0.83Al0.17As/ In0.83Ga0.17As thin films which also alternate with each other but in a reverse order, and the thickness ratios of In0.83Al0.17As/In0.83Ga0.17As in each group are 1:9, 2:8, …, finally coming to …, 8:2, 9:1, respectively. It is similar to the chirped In0.53Ga0.47As/In0.52Al0.48As superlattice-graded bandgap layer used in some literatures [21–23].

**Figure 4(a)** shows the two different growth profiles of the In composition of In<sup>x</sup> Al1−xAs buffer for the InP-based In0.83Ga0.17As PDs. Through a comprehensive comparison, sample S1 with continuously graded Inx Al1−xAs MBL shows better properties on both material and devices than that of sample S2 with step-graded Inx Al1−xAs MBL. As shown in the cross-sectional transmission electron micrographs (XTEMs) of **Figure 5**, almost none evident TD was found in the absorption layer of sample S1. By contrast, some obvious TDs exist in the absorption layer of sample S2. Unsurprisingly, sample S1 shows superior device performance compared

to sample S2. By measuring the mesa-type devices in diameter size of 200 μm, the room temperature (RT) dark current of the sample S1 was smaller than that of sample S2 at low reverse bias condition, as shown in **Figure 4(b)**. At reverse bias of −10 mV, the dark currents are 259 nA

0 500 1000 1500 2000

Sample S1 Sample S2

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Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch

Sample S1 Sample S2

Thickness (nm)

0.0 0.2 0.4 0.6 0.8 1.0

Reverse Bias (V)

**Figure 4.** (a) The profiles of the indium composition vs. the grown thickness, and (b) I-V characteristics at 300 K of

0.5

1E-7

1E-6

Dark Current (A)

samples S1 and S2.

**Figure 5.** XTEM images of samples S1 and S2.

0.6

In content

0.7

0.8

(a)

(b)

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch http://dx.doi.org/10.5772/intechopen.70259 209

**Figure 4.** (a) The profiles of the indium composition vs. the grown thickness, and (b) I-V characteristics at 300 K of samples S1 and S2.

**Figure 5.** XTEM images of samples S1 and S2.

InGaAs and InAlAs at the interfaces, digital-graded superlattices (DGSLs) were grown at the interfaces, respectively. First, DGSL1 consisted of nine groups of In0.83Ga0.17As/In0.83Al0.17As thin films which alternate with each other. The thickness of each group is about 8 nm. The thickness ratios of In0.83Ga0.17As/In0.83Al0.17As in each group are 1:9, 2:8, …, finally coming to …, 8:2, 9:1, respectively. Similarly, DGSL2 was comprised of nine groups of In0.83Al0.17As/ In0.83Ga0.17As thin films which also alternate with each other but in a reverse order, and the thickness ratios of In0.83Al0.17As/In0.83Ga0.17As in each group are 1:9, 2:8, …, finally coming to …, 8:2, 9:1, respectively. It is similar to the chirped In0.53Ga0.47As/In0.52Al0.48As superlattice-graded

for the InP-based In0.83Ga0.17As PDs. Through a comprehensive comparison, sample S1 with

transmission electron micrographs (XTEMs) of **Figure 5**, almost none evident TD was found in the absorption layer of sample S1. By contrast, some obvious TDs exist in the absorption layer of sample S2. Unsurprisingly, sample S1 shows superior device performance compared

Al1−xAs MBL shows better properties on both material and devices

Al1−xAs MBL. As shown in the cross-sectional

Al1−xAs buffer

DGSL1

DGSL2

**Figure 4(a)** shows the two different growth profiles of the In composition of In<sup>x</sup>

bandgap layer used in some literatures [21–23].

S.I. InP substrate

**Figure 3.** The schematic structure of In0.83Ga0.17As photodetectors.

200 nm N+ InP buffer 100 nm N+ In0.52Al0.48As

1.9 µm N+ linearly or step graded InxAl1-xAs buffer x: 0.52→0.83

n- 72 nm DGSL1

1.5 µm n- In0.83Ga0.17As absorption layer

72 nm undoped DGSL2 0.53 µm p+ In0.83Al0.17As cap layer

208 Epitaxy

than that of sample S2 with step-graded Inx

continuously graded Inx

to sample S2. By measuring the mesa-type devices in diameter size of 200 μm, the room temperature (RT) dark current of the sample S1 was smaller than that of sample S2 at low reverse bias condition, as shown in **Figure 4(b)**. At reverse bias of −10 mV, the dark currents are 259 nA (8.25 × 10−4 A/cm<sup>2</sup> ) and 473 nA (1.51 × 10−3 A/cm<sup>2</sup> ) for samples S1 and S2, respectively. Therefore, continuously graded Inx Al1−xAs MBL scheme is more suitable for the growth of In0.83Ga0.17As PDs on InP substrate.

In the previous continuously graded InAlAs MBLs, all the In contents were graded from 0.52. The initial lattice constant matches to the InP substrate. Strain energies in the continuously graded buffer increase gradually with increasing In content. Then, dispersed misfit dislocations will be generated to release the strain energy slowly in the buffer. Little overlap and interaction would occur among the TDs. Thus, an enough thick buffer is needed to slide and annihilate the TDs during the growth process. Here, we propose another method to promote the overlap of dispersed TDs by introducing an abrupt initial In<sup>y</sup> Al1−yAs (y > 0.52) layer of InAlAs MBL on InP substrate. This abrupt initial layer, mismatched with the InP substrate, is similar to a step buffer. But it was combined with a continuously graded In<sup>x</sup> Al1−xAs buffer with x graded from y to the composition needed. This buffer scheme, so-called one-step continuously graded buffer, is expected to possess the advantages of both step and continuously graded schemes. The strain is released rapidly by introducing the initial abrupt buffer, and TDs were restricted between the step buffer and the InP substrate. We show that this method results in lower TDD in the final-graded buffer and the InGaAs absorption layer in the wavelength-extended PD structure. This buffer strategy has been also adopted and demonstrated for the growth of InAsx P1−x on InP by MOCVD method [23].

To investigate the effect of this buffer scheme on the properties of InP-based In0.83Ga0.17As PDs with a lattice mismatch of 2.1%. Samples with three different initial In compositions were grown by GSMBE. As a comparison, the In composition grading profile and growth method of sample E were remained the same as before [19, 24]. For the deposition of wafers with onestep continuously graded Iny Al1−yAs buffer, the In contents y of the one-step buffer were set to be 0.68 and 0.77 with the thicknesses of 0.5 and 0.7 μm, respectively. Then, the one-step buffer was followed by a continuously graded Inx Al1−xAs buffer with In composition graded from 0.68 or 0.77 to 0.83, respectively. The growth temperatures were graded from 520 to 490°C and 505 to 490°C, respectively. They were defined as samples F and G, respectively. Since the total thickness of the InAlAs MBL was maintained to be 1.9 μm, the higher the initial indium content, the lower the mismatch grading rate of the continuously graded Iny Al1−yAs buffer, as shown in **Figure 6(a)**. The mismatch grading rates for samples E, F, and G were 1.08, 0.702, and 0.312% μm−1, respectively.

epilayer. The strain gradient is proportional to the compositional grading rate in a linearly graded MBL. In this section, 100-nm thin In0.8Ga0.2As layers were grown atop on the 1.6 μm

**Figure 6.** (a) In content vs. the grown thickness, (b) RT PL spectra, and (c) measured I-V characteristics of samples E, F,

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch

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211

and B, respectively. Since the lattice mismatch between the In0.8Ga0.2As layer and the InP substrate is about +1.86%. The strain gradients correspond to be around +1.2% μm−1 and +3.1% μm−1, for samples A and B, respectively. The In composition profiles versus growth thickness for the two samples were shown in **Figure 7**. The effects of strain gradient on the surface, structural, and optical properties of the In0.80Ga0.20As metamorphic structures were

The results show that although lower strain gradient caused a slightly rougher surface of the structure shown in **Figure 8**, but it also brought about a larger degree of lattice relaxation and a less residual strain in the top In0.8Ga0.2As layer. **Figure 9** shows the measurement of highresolution XRD (HRXRD) reciprocal space mapping (RSM). Some results were extracted from

The peaks of the top In0.80Ga0.20As layer of sample A are stronger than that of sample B in both (0 0 4) and (2 2 4) reflections, indicating higher lattice quality of sample A. The lattice relaxation degree is 82.5% for sample A and 77.8% for sample B, respectively. Correspondingly, the residual strain of sample A is smaller than that of sample B. Therefore, it can be concluded that the use of lower strain gradient in the linearly graded MBL is beneficial to the lattice relaxation and the release of residual strain. In addition, a nearly fully strained thin top layer of the structure was observed, indicating a two-step relaxation procedure of the linearly graded MBL as predicted by the Tersoff's model [26]. However, it is still not suffi-

cient to achieve a full relaxation even in the linearly graded Inx

Ga1−xAs MBLs on InP, which were defined as samples A

Ga1−xAs MBL with a mismatch

and 0.6 μm linearly graded In<sup>x</sup>

and G. Reprinted with permission from Elsevier.

the RSM and listed in **Table 2**.

investigated [25].

By using this buffer scheme, it was found that RT PL intensity of In0.83Ga0.17As layer was enhanced in sequence from E to G as shown in **Figure 6(b)**, indicating a promoted optical quality of the absorption layer. For the mesa type devices in the diameter size of 200 μm, the dark current was deceased along with the increase of initial In content y from 0.52 to 0.68 and 0.77 under low reverse bias condition. At a reverse bias of −10 mV, the dark current halves for the sample G with respect to that of the sample E at 300 K, as shown in **Figure 6(c)**. This validated the suppression effects of one-step continuously graded In<sup>y</sup> Al1−yAs MBL on the TDs.

#### **2.3. Strain gradient and composition overshoot in the metamorphic buffer layer**

An ideal buffer cannot only restrict the misfit dislocations in the inactive region and reduce the TDD in the absorption layer but also immensely reduce the residual lattice strain in the

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch http://dx.doi.org/10.5772/intechopen.70259 211

(8.25 × 10−4 A/cm<sup>2</sup>

210 Epitaxy

continuously graded Inx

PDs on InP substrate.

for the growth of InAsx

step continuously graded Iny

and 0.312% μm−1, respectively.

was followed by a continuously graded Inx

) and 473 nA (1.51 × 10−3 A/cm<sup>2</sup>

the overlap of dispersed TDs by introducing an abrupt initial In<sup>y</sup>

is similar to a step buffer. But it was combined with a continuously graded In<sup>x</sup>

P1−x on InP by MOCVD method [23].

content, the lower the mismatch grading rate of the continuously graded Iny

dated the suppression effects of one-step continuously graded In<sup>y</sup>

**2.3. Strain gradient and composition overshoot in the metamorphic buffer layer**

) for samples S1 and S2, respectively. Therefore,

Al1−yAs (y > 0.52) layer of

Al1−xAs buffer

Al1−yAs buffer, as

Al1−yAs MBL on the TDs.

Al1−xAs MBL scheme is more suitable for the growth of In0.83Ga0.17As

Al1−yAs buffer, the In contents y of the one-step buffer were set to

Al1−xAs buffer with In composition graded from

In the previous continuously graded InAlAs MBLs, all the In contents were graded from 0.52. The initial lattice constant matches to the InP substrate. Strain energies in the continuously graded buffer increase gradually with increasing In content. Then, dispersed misfit dislocations will be generated to release the strain energy slowly in the buffer. Little overlap and interaction would occur among the TDs. Thus, an enough thick buffer is needed to slide and annihilate the TDs during the growth process. Here, we propose another method to promote

InAlAs MBL on InP substrate. This abrupt initial layer, mismatched with the InP substrate,

with x graded from y to the composition needed. This buffer scheme, so-called one-step continuously graded buffer, is expected to possess the advantages of both step and continuously graded schemes. The strain is released rapidly by introducing the initial abrupt buffer, and TDs were restricted between the step buffer and the InP substrate. We show that this method results in lower TDD in the final-graded buffer and the InGaAs absorption layer in the wavelength-extended PD structure. This buffer strategy has been also adopted and demonstrated

To investigate the effect of this buffer scheme on the properties of InP-based In0.83Ga0.17As PDs with a lattice mismatch of 2.1%. Samples with three different initial In compositions were grown by GSMBE. As a comparison, the In composition grading profile and growth method of sample E were remained the same as before [19, 24]. For the deposition of wafers with one-

be 0.68 and 0.77 with the thicknesses of 0.5 and 0.7 μm, respectively. Then, the one-step buffer

0.68 or 0.77 to 0.83, respectively. The growth temperatures were graded from 520 to 490°C and 505 to 490°C, respectively. They were defined as samples F and G, respectively. Since the total thickness of the InAlAs MBL was maintained to be 1.9 μm, the higher the initial indium

shown in **Figure 6(a)**. The mismatch grading rates for samples E, F, and G were 1.08, 0.702,

By using this buffer scheme, it was found that RT PL intensity of In0.83Ga0.17As layer was enhanced in sequence from E to G as shown in **Figure 6(b)**, indicating a promoted optical quality of the absorption layer. For the mesa type devices in the diameter size of 200 μm, the dark current was deceased along with the increase of initial In content y from 0.52 to 0.68 and 0.77 under low reverse bias condition. At a reverse bias of −10 mV, the dark current halves for the sample G with respect to that of the sample E at 300 K, as shown in **Figure 6(c)**. This vali-

An ideal buffer cannot only restrict the misfit dislocations in the inactive region and reduce the TDD in the absorption layer but also immensely reduce the residual lattice strain in the

**Figure 6.** (a) In content vs. the grown thickness, (b) RT PL spectra, and (c) measured I-V characteristics of samples E, F, and G. Reprinted with permission from Elsevier.

epilayer. The strain gradient is proportional to the compositional grading rate in a linearly graded MBL. In this section, 100-nm thin In0.8Ga0.2As layers were grown atop on the 1.6 μm and 0.6 μm linearly graded In<sup>x</sup> Ga1−xAs MBLs on InP, which were defined as samples A and B, respectively. Since the lattice mismatch between the In0.8Ga0.2As layer and the InP substrate is about +1.86%. The strain gradients correspond to be around +1.2% μm−1 and +3.1% μm−1, for samples A and B, respectively. The In composition profiles versus growth thickness for the two samples were shown in **Figure 7**. The effects of strain gradient on the surface, structural, and optical properties of the In0.80Ga0.20As metamorphic structures were investigated [25].

The results show that although lower strain gradient caused a slightly rougher surface of the structure shown in **Figure 8**, but it also brought about a larger degree of lattice relaxation and a less residual strain in the top In0.8Ga0.2As layer. **Figure 9** shows the measurement of highresolution XRD (HRXRD) reciprocal space mapping (RSM). Some results were extracted from the RSM and listed in **Table 2**.

The peaks of the top In0.80Ga0.20As layer of sample A are stronger than that of sample B in both (0 0 4) and (2 2 4) reflections, indicating higher lattice quality of sample A. The lattice relaxation degree is 82.5% for sample A and 77.8% for sample B, respectively. Correspondingly, the residual strain of sample A is smaller than that of sample B. Therefore, it can be concluded that the use of lower strain gradient in the linearly graded MBL is beneficial to the lattice relaxation and the release of residual strain. In addition, a nearly fully strained thin top layer of the structure was observed, indicating a two-step relaxation procedure of the linearly graded MBL as predicted by the Tersoff's model [26]. However, it is still not sufficient to achieve a full relaxation even in the linearly graded Inx Ga1−xAs MBL with a mismatch

**Figure 7.** The indium composition versus growth thickness for the two samples. Reprinted with permission from IOP.

**Figure 8.** AFM images of (a) sample A and (b) sample B. Reprinted with permission from IOP.

grading rate of 1.2% μm−1 in our experiments. Therefore, to further increase the lattice relaxation degree, composition "overshoot" in the MBL should be introduced.

It was proposed that a dislocation-free portion will be formed when the thickness of the buffer is excess of a value of in the linearly graded buffer. The value of *Zc* can be expressed as [27]:

$$Z\_{\varepsilon} = \mathcal{W} \text{ - (2\lambda\text{lbc }\varepsilon\text{ })^\vee }\tag{1}$$

**Figure 9.** HRXRD RSMs of (a) sample A and (b) sample B. Reprinted with permission from IOP.

**Relaxation degree (%)**

**Table 2.** Results extracted from the RSM measurements (reprinted with permission from IOP).

**Parallel mismatch (%)**

M 1.907 0.806 82.5 1.573 2.322 −3.27 −4.4 N 1.885 0.803 77.8 1.466 2.364 −4.11 −7.0

**Perpendicular mismatch (%)**

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213

**Residual strain (10−3) XRD RSM Tersoff's** 

**model**

**Indium content**

**Samples Bulk mismatch (%)**

where *W* is the total thickness of the buffer, *λ* is the energy per unit length of the dislocation, *c* is the appropriate elastic constant for biaxial strain, *b* is the misfit component of the Burgers vector of dislocation, and *ε*′ is the strain gradient. The residual strain in this dislocation-free portion can be given by:

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch http://dx.doi.org/10.5772/intechopen.70259 213

**Figure 9.** HRXRD RSMs of (a) sample A and (b) sample B. Reprinted with permission from IOP.

grading rate of 1.2% μm−1 in our experiments. Therefore, to further increase the lattice relax-

**Figure 7.** The indium composition versus growth thickness for the two samples. Reprinted with permission from IOP.

It was proposed that a dislocation-free portion will be formed when the thickness of the buffer

where *W* is the total thickness of the buffer, *λ* is the energy per unit length of the dislocation, *c* is the appropriate elastic constant for biaxial strain, *b* is the misfit component of the Burgers vector of dislocation, and *ε*′ is the strain gradient. The residual strain in this dislocation-free

can be expressed as [27]:

) 1⁄<sup>2</sup> (1)

ation degree, composition "overshoot" in the MBL should be introduced.

**Figure 8.** AFM images of (a) sample A and (b) sample B. Reprinted with permission from IOP.

is excess of a value of in the linearly graded buffer. The value of *Zc*

*Zc* = *W* - (2*λ*/*bc ε*′

portion can be given by:

212 Epitaxy


**Table 2.** Results extracted from the RSM measurements (reprinted with permission from IOP).

$$
\overline{\varepsilon} = \mathcal{W}\varepsilon' - Z\_{\varepsilon}\varepsilon' = \langle \Sigma\lambda \ \varepsilon' \mathcal{W}\varepsilon \rangle^{\omega} \tag{2}
$$

**2.4. Adoption of DA for dislocation restriction in the metamorphic buffer layer**

**Figure 11.** AFM images of (a) sample 1, (b) sample 2, (c) sample 3, and (d) sample 4.

*i*=0

are so thin that they would intermix with each other significantly.

*n* is the total number of thin layers and m is the composition in each thin layer. These layers

*n di* × *m* \_\_\_\_\_

of the lattice-mismatched system [29].

) through the equation:

∑

each period (*di*

Recently, it was confirmed that the insertion of InAs/In0.52Al0.48As DA intermediate layer in the InAlAs MBL can restrain the TDs effectively and improve the structural and optical qualities

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch

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215

By alternating the epitaxy of two or more thin layers, DA intermediate layers can be formed. They are expected to reduce the strain energy, restrain the three dimensional (3D) growth, and increase the critical thickness. If the total thickness of one period is set to be *d*, the needed composition *z* of the DA layer can be achieved by adjusting the thickness of each thin layer in

*<sup>d</sup>* <sup>=</sup> *<sup>z</sup>* (3)

Therefore, in the linearly graded buffer, an "overshoot" of the mismatch by an amount is needed to make the graded buffer be lattice matched to the following layers. Therefore, to judge this hypothesis, two sets of In0.78Ga0.22As/In0.78Al0.22As quantum wells (QWs) and In0.84Ga0.16As PD structures with and without compositional overshoot in the linearly graded InAlAs MBLs were grown. The effects on the material properties and the lattice relaxation degree have been investigated in detail [28].

As shown in **Figure 10**, the QW structures consist of a 100-nm-thick InP buffer, a 1.7-μm-thick linearly graded Inx Al1−xAs MBL, and a three periods of In0.78Ga0.22As/In0.78Al0.22As QWs. The thicknesses of well and barrier layers are 10 and 12 nm respectively, as shown in **Figure 10(a)** and **(b)**. The In composition x in the In<sup>x</sup> Al1−xAs buffer of this structure was graded from 0.52 to 0.78 for sample 1 and from 0.52 to 0.82 for sample 2. For the In0.84Ga0.16As PD structures, the In composition x in the Inx Al1−xAs buffer was designed grading from 0.52 to 0.84 for sample 3 and from 0.52 to 0.88 for sample 4, as shown in **Figure 10(c)** and **(d)**, respectively. Therefore, compositional overshoot of about 0.04 was introduced in the end of the In<sup>x</sup> Al1−xAs MBLs of samples 2 and 4.

From the AFM, PL, and HRXRD RSM measurement results shown in **Figure 11**, **Figure 12**, and **Table 3**, respectively, we have concluded that the use of compositional overshoot in the InAlAs MBLs not only reduces the surface roughness but also enhances the optical quality and the lattice relaxation degree for both QW and PD structures.

**Figure 10.** The indium composition profiles in the InAlAs buffer and schematic structures of samples (a) 1, (b) 2, (c) 3, and (d) 4, respectively.

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch http://dx.doi.org/10.5772/intechopen.70259 215

**Figure 11.** AFM images of (a) sample 1, (b) sample 2, (c) sample 3, and (d) sample 4.

¯¯ε = *W ε*′ - *Zc ε*′ = (2*λ ε*′

degree have been investigated in detail [28].

and **(b)**. The In composition x in the In<sup>x</sup>

linearly graded Inx

214 Epitaxy

samples 2 and 4.

and (d) 4, respectively.

In composition x in the Inx

Therefore, in the linearly graded buffer, an "overshoot" of the mismatch by an amount is needed to make the graded buffer be lattice matched to the following layers. Therefore, to judge this hypothesis, two sets of In0.78Ga0.22As/In0.78Al0.22As quantum wells (QWs) and In0.84Ga0.16As PD structures with and without compositional overshoot in the linearly graded InAlAs MBLs were grown. The effects on the material properties and the lattice relaxation

As shown in **Figure 10**, the QW structures consist of a 100-nm-thick InP buffer, a 1.7-μm-thick

thicknesses of well and barrier layers are 10 and 12 nm respectively, as shown in **Figure 10(a)**

to 0.78 for sample 1 and from 0.52 to 0.82 for sample 2. For the In0.84Ga0.16As PD structures, the

and from 0.52 to 0.88 for sample 4, as shown in **Figure 10(c)** and **(d)**, respectively. Therefore,

From the AFM, PL, and HRXRD RSM measurement results shown in **Figure 11**, **Figure 12**, and **Table 3**, respectively, we have concluded that the use of compositional overshoot in the InAlAs MBLs not only reduces the surface roughness but also enhances the optical quality

**Figure 10.** The indium composition profiles in the InAlAs buffer and schematic structures of samples (a) 1, (b) 2, (c) 3,

compositional overshoot of about 0.04 was introduced in the end of the In<sup>x</sup>

and the lattice relaxation degree for both QW and PD structures.

Al1−xAs MBL, and a three periods of In0.78Ga0.22As/In0.78Al0.22As QWs. The

Al1−xAs buffer was designed grading from 0.52 to 0.84 for sample 3

Al1−xAs buffer of this structure was graded from 0.52

/*bc*) 1⁄<sup>2</sup> (2)

Al1−xAs MBLs of

#### **2.4. Adoption of DA for dislocation restriction in the metamorphic buffer layer**

Recently, it was confirmed that the insertion of InAs/In0.52Al0.48As DA intermediate layer in the InAlAs MBL can restrain the TDs effectively and improve the structural and optical qualities of the lattice-mismatched system [29].

By alternating the epitaxy of two or more thin layers, DA intermediate layers can be formed. They are expected to reduce the strain energy, restrain the three dimensional (3D) growth, and increase the critical thickness. If the total thickness of one period is set to be *d*, the needed composition *z* of the DA layer can be achieved by adjusting the thickness of each thin layer in each period (*di* ) through the equation:

$$\sum\_{i=0}^{n} \frac{d\_i \times m}{d} = \mathbf{z} \tag{3}$$

*n* is the total number of thin layers and m is the composition in each thin layer. These layers are so thin that they would intermix with each other significantly.

**Figure 12.** The RT PL spectra of (a) sample 1, (b) sample 2, (c) sample 3, and (d) sample 4.


500 nm In<sup>x</sup>

from the equation:

(*d*2

Al1−xAs buffer. The In composition of the intermediate layer, y, was designed to be

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217

) and In0.52Al0.48As

the same as the composition in the inserted positions of the buffer, i.e., y = 0.62 and 0.71 for DA intermediate layers I and II, respectively. The DA layers were formed by the alternating epitaxy of very thin InAs and In0.52Al0.48As layers. The thickness *d* of a period including one

**Figure 13.** (a) Schematic design structure and (b) XTEM images of the samples. Reprinted with permission from IOP.

) in each period were adjusted to achieve the equivalent In composition y and obtained

InAs layer and one In0.52Al0.48As layer was 1 nm. The thicknesses of InAs (*d*<sup>1</sup>

**Table 3.** Results extracted from RSM measurements of samples 1–4.

As shown schematically in **Figure 13(a)**, two In0.8Al0.2As/In0.8Ga0.2As QW structures were grown on linearly graded InAlAs MBL on InP. As a reference, the In composition x in sample A was continuously graded from 0.52 to 0.8. For sample B, two In<sup>y</sup> Al1−yAs DA intermediate layers of about 100 nm were inserted into the graded InAlAs MBL, with a separation of every Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch http://dx.doi.org/10.5772/intechopen.70259 217

**Figure 13.** (a) Schematic design structure and (b) XTEM images of the samples. Reprinted with permission from IOP.

500 nm In<sup>x</sup> Al1−xAs buffer. The In composition of the intermediate layer, y, was designed to be the same as the composition in the inserted positions of the buffer, i.e., y = 0.62 and 0.71 for DA intermediate layers I and II, respectively. The DA layers were formed by the alternating epitaxy of very thin InAs and In0.52Al0.48As layers. The thickness *d* of a period including one InAs layer and one In0.52Al0.48As layer was 1 nm. The thicknesses of InAs (*d*<sup>1</sup> ) and In0.52Al0.48As (*d*2 ) in each period were adjusted to achieve the equivalent In composition y and obtained from the equation:

As shown schematically in **Figure 13(a)**, two In0.8Al0.2As/In0.8Ga0.2As QW structures were grown on linearly graded InAlAs MBL on InP. As a reference, the In composition x in sample

layers of about 100 nm were inserted into the graded InAlAs MBL, with a separation of every

Al1−yAs DA intermediate

**Residual strain** 

**(10−3)**

A was continuously graded from 0.52 to 0.8. For sample B, two In<sup>y</sup>

**Figure 12.** The RT PL spectra of (a) sample 1, (b) sample 2, (c) sample 3, and (d) sample 4.

**Parallel mismatch (%)**

 0.774 1.675 1.459 1.510 87.2 −2.13 0.781 1.730 1.716 2.006 96.5 −0.60 0.838 2.125 2.069 2.324 97.4 −0.54 0.840 2.136 2.095 2.288 99.4 −0.45

**Perpendicular mismatch (%)**

**Relax. Degree (%)**

**Samples Indium content Cubic mismatch** 

216 Epitaxy

**(%)**

**Table 3.** Results extracted from RSM measurements of samples 1–4.

$$\begin{cases} \frac{d\_1 + d\_2 \times 0.52}{d} = y\_2\\ \quad d\_1 + d\_2 = d\_1 \end{cases}$$

*d*1 and *d*<sup>2</sup> are so small that a considerable material intermixing is expected to occur between the two layers since that the thinnest layer is a submonolayer, and the thickest layer is only about two monolayers (MLs). After the growth of the MBL, three 10-nm In0.8Ga0.2As QWs sandwiched by In0.8Al0.2As barriers were grown. The thickness was 12 nm for the In0.8Al0.2As barrier between the two wells and 100 nm for the first and last barriers.

The measurement results were shown in Figures **13(b)–16**. Although the insertion of InAs/ In0.52Al0.48As DA intermediate layers caused a slightly decrease of lattice relaxation degree listed in **Table 4**, the TDD has been decreased immensely in the following InAlAs MBLs as shown in the XTEM images of **Figure 13(b)**, and both the PL and AFM RMS have been improved markedly, indicating the positive effect of the DAs on the structural and optical qualities. However, to maximize these effects, the composition, number, and thickness of the DA should be further optimized.

**Figure 16.** HRXRD RSMs of (a) sample A and (b) sample B. Reprinted with permission from IOP.

AFM surface images of (a) sample A and (b) sample B. Reprinted with permission from IOP.

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch

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219

**Figure 15.** 20 × 20 μm<sup>2</sup>

**Figure 14.** PL spectra of the samples at 300 and 77 K. Reprinted with permission from IOP.

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch http://dx.doi.org/10.5772/intechopen.70259 219

{

DA should be further optimized.

*d*1 and *d*<sup>2</sup>

218 Epitaxy

*<sup>d</sup>*<sup>1</sup> <sup>+</sup> *<sup>d</sup>*<sup>2</sup> <sup>×</sup> 0.52 \_\_\_\_\_\_\_\_\_

*<sup>d</sup>* <sup>=</sup> *<sup>y</sup>*

are so small that a considerable material intermixing is expected to occur between

*d*<sup>1</sup> + *d*<sup>2</sup> = *d*

the two layers since that the thinnest layer is a submonolayer, and the thickest layer is only about two monolayers (MLs). After the growth of the MBL, three 10-nm In0.8Ga0.2As QWs sandwiched by In0.8Al0.2As barriers were grown. The thickness was 12 nm for the In0.8Al0.2As

The measurement results were shown in Figures **13(b)–16**. Although the insertion of InAs/ In0.52Al0.48As DA intermediate layers caused a slightly decrease of lattice relaxation degree listed in **Table 4**, the TDD has been decreased immensely in the following InAlAs MBLs as shown in the XTEM images of **Figure 13(b)**, and both the PL and AFM RMS have been improved markedly, indicating the positive effect of the DAs on the structural and optical qualities. However, to maximize these effects, the composition, number, and thickness of the

barrier between the two wells and 100 nm for the first and last barriers.

**Figure 14.** PL spectra of the samples at 300 and 77 K. Reprinted with permission from IOP.

**Figure 15.** 20 × 20 μm<sup>2</sup> AFM surface images of (a) sample A and (b) sample B. Reprinted with permission from IOP.

**Figure 16.** HRXRD RSMs of (a) sample A and (b) sample B. Reprinted with permission from IOP.


**Table 4.** Results of XRD RSM measurements (reprinted with permission from IOP).

## **3. Trials to move In0.83Ga0.17As PD from InP to GaAs substrate**

#### **3.1. Comparison of GaAs- and InP-based In0.83Ga0.17As PDs with different lattice mismatches**

Currently, most of the wavelength-extended InGaAs PD structures were grown on the InP substrate [2]. Even if a larger lattice mismatch will be introduced, GaAs may still be an attractive substrate for fabrication of InGaAs PDs with large size epitaxial wafers as well as FPAs with more pixels for the advantages of robustness, lower cost, and larger size. However, there are few attempts to transfer In<sup>x</sup> Ga1−xAs (x > 0.53) PDs from InP to GaAs substrate.

though high lattice relaxation degrees were acquired in both structures, a relatively higher residual strain was observed in the GaAs-based In0.83Ga0.17As PD structure than that in the

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch

thin thickness could not well accommodate the relatively higher lattice mismatch between

From XRD measurement shown in **Figure 18(a)**, the intensity ratio of epilayer/substrate for sample S is much smaller than that of sample P. Moreover, the FWHM value of the In0.83Ga0.17As

**Figure 18.** (a) High-resolution (0 0 4) *ω* − 2*θ* XRD scan curves, and (b) temperature-dependent reverse I-V characteristics

of the two samples. Reprinted with permission from Elsevier.

Al1−xAs MBL with relatively

. Reprinted with permission from

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221

InP-based structure. This implies that the continuously graded In<sup>x</sup>

**Figure 17.** AFM images of (a) sample S, and (b) sample P. The scan area is 20 × 20 μm<sup>2</sup>

In0.83Ga0.17As and GaAs.

Elsevier.

In our recent researches, two In0.83Ga0.17As PD structures with continuously graded InAlAs MBL were grown on (1 0 0)-oriented S-doped InP or GaAs epi-ready substrates by GSMBE [30]. In this way, the lattice mismatch will increase from +2.1 to +5.9% for In0.83Ga0.17As on GaAs compared to that on InP. The growth condition of InP-based In0.83Ga0.17As PD structure was the same as that in the previous study. For the deposition of the GaAs-based In0.83Ga0.17As PD structure, the InAlAs MBL started with a 0.1-μm-thick highly Si-doped N+ In0.1Al0.9As layer grown at 530°C, followed by a 1.9-μm-thick compositionally graded In<sup>x</sup> Al1−xAs layer with In composition x graded from 0.1 to 0.87, and the substrate temperature graded from 530 to 460°C. The mismatch grading rate of 3.1% μm−1 for sample A was larger than that of 1.1% μm−1 for sample B. Then, a 0.65-μm In0.83Al0.17As template was grown at 460°C in the end of the graded buffer. The growth temperatures of the InGaAs absorption layer and the InAlAs cap layer were 490°C, and the doping level in each layer of both structures was kept uniform. They were renamed as samples S and P for GaAs-based and InP-based In0.83Ga0.17As PD structures, respectively. Their features have been evaluated on both material qualities and device performances.

As the AFM images shows in **Figure 17**, typical anisotropic features of the surface are detected along the [1 1 0] or [1 −1 0] direction in both samples. While the oval-like defects on the surface of sample S seem more distinct, and the undulations of the pattern are larger than those of sample P. The RMS roughness values are 11.3 nm and 4.8 nm for samples S and P, respectively, as shown in **Figure 17**. The relatively smaller RMS value of surface roughness indicates a better crystalline quality of InP-based In0.83Ga0.17As PD structure. The results of measured symmetric (0 0 4) and asymmetric (2 2 4) reflection RSMs (not shown here) showed that Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch http://dx.doi.org/10.5772/intechopen.70259 221

**Figure 17.** AFM images of (a) sample S, and (b) sample P. The scan area is 20 × 20 μm<sup>2</sup> . Reprinted with permission from Elsevier.

**3. Trials to move In0.83Ga0.17As PD from InP to GaAs substrate**

**Table 4.** Results of XRD RSM measurements (reprinted with permission from IOP).

A QW structure 0.760 82.7 −3.5

B DA II 0.653 67.1 −2.9

**mismatches**

220 Epitaxy

are few attempts to transfer In<sup>x</sup>

device performances.

**3.1. Comparison of GaAs- and InP-based In0.83Ga0.17As PDs with different lattice** 

PD structure, the InAlAs MBL started with a 0.1-μm-thick highly Si-doped N+

layer grown at 530°C, followed by a 1.9-μm-thick compositionally graded In<sup>x</sup>

Currently, most of the wavelength-extended InGaAs PD structures were grown on the InP substrate [2]. Even if a larger lattice mismatch will be introduced, GaAs may still be an attractive substrate for fabrication of InGaAs PDs with large size epitaxial wafers as well as FPAs with more pixels for the advantages of robustness, lower cost, and larger size. However, there

**Sample Epitaxy layers Indium content (y) Degree of relaxation, R (%) Residual strain, ε (10−3)**

QW structure 0.736 70.2 −4.3

DA I 0.580 54.9 −1.8

In our recent researches, two In0.83Ga0.17As PD structures with continuously graded InAlAs MBL were grown on (1 0 0)-oriented S-doped InP or GaAs epi-ready substrates by GSMBE [30]. In this way, the lattice mismatch will increase from +2.1 to +5.9% for In0.83Ga0.17As on GaAs compared to that on InP. The growth condition of InP-based In0.83Ga0.17As PD structure was the same as that in the previous study. For the deposition of the GaAs-based In0.83Ga0.17As

with In composition x graded from 0.1 to 0.87, and the substrate temperature graded from 530 to 460°C. The mismatch grading rate of 3.1% μm−1 for sample A was larger than that of 1.1% μm−1 for sample B. Then, a 0.65-μm In0.83Al0.17As template was grown at 460°C in the end of the graded buffer. The growth temperatures of the InGaAs absorption layer and the InAlAs cap layer were 490°C, and the doping level in each layer of both structures was kept uniform. They were renamed as samples S and P for GaAs-based and InP-based In0.83Ga0.17As PD structures, respectively. Their features have been evaluated on both material qualities and

As the AFM images shows in **Figure 17**, typical anisotropic features of the surface are detected along the [1 1 0] or [1 −1 0] direction in both samples. While the oval-like defects on the surface of sample S seem more distinct, and the undulations of the pattern are larger than those of sample P. The RMS roughness values are 11.3 nm and 4.8 nm for samples S and P, respectively, as shown in **Figure 17**. The relatively smaller RMS value of surface roughness indicates a better crystalline quality of InP-based In0.83Ga0.17As PD structure. The results of measured symmetric (0 0 4) and asymmetric (2 2 4) reflection RSMs (not shown here) showed that

Ga1−xAs (x > 0.53) PDs from InP to GaAs substrate.

In0.1Al0.9As

Al1−xAs layer

though high lattice relaxation degrees were acquired in both structures, a relatively higher residual strain was observed in the GaAs-based In0.83Ga0.17As PD structure than that in the InP-based structure. This implies that the continuously graded In<sup>x</sup> Al1−xAs MBL with relatively thin thickness could not well accommodate the relatively higher lattice mismatch between In0.83Ga0.17As and GaAs.

From XRD measurement shown in **Figure 18(a)**, the intensity ratio of epilayer/substrate for sample S is much smaller than that of sample P. Moreover, the FWHM value of the In0.83Ga0.17As

**Figure 18.** (a) High-resolution (0 0 4) *ω* − 2*θ* XRD scan curves, and (b) temperature-dependent reverse I-V characteristics of the two samples. Reprinted with permission from Elsevier.

layer of sample S is much larger than that of sample P with values of 1638 and 644 arcsec, respectively. These suggests that the crystal quality of the In0.83Ga0.17As grown on the GaAs substrate is not so good as that grown on the InP substrate with a higher lattice mismatch by using the same type of MBL. This verified one of the conclusions in our previous work [31] that full relaxation and favorable optical property of the InGaAs layer could not occur for the PD wafers with a mismatch grading rate of 2.4% μm−1 or above but occur for the wafers with a mismatch grading rate of about 1.2% μm−1 or lower.

Correspondingly, quite different temperature dependent I-V characteristics of the two PD chips with diameter of 300 μm were observed in the temperature range from 77 to 300 K, as shown in **Figure 18(b)**. At reverse bias of −10 mV, the dark currents of 2.28 μA at 300 K and 2.17 nA at 77 K for sample S are much larger than that of 674 nA at 300 K and 3.99 pA at 77 K for sample P, respectively, while the zero bias R0 A of the GaAs-based PD is comparable with that of the InP-based PD at 300 K [32], as shown in **Table 5**. This indicates that via appropriate structural design, GaAs may act as a feasible substrate for replacement of InP for In0.83Ga0.17As PDs in some low-end application area around RT.

To analyze the cause of the different electrical performances of the two samples, XTEM measurement was performed firstly, as shown in **Figure 19**. From the XTEM bright-field images, it is obvious that the majority of misfit dislocations of sample P are mainly localized at the early stage of the continuously graded Inx Al1−xAs MBL. Threading dislocations have been prevented from propagating into the In0.83Ga0.17As absorption layer by the continuously graded InAlAs MBL on InP under this strain gradient condition. However, it seems that the effect of the continuously graded InAlAs MBL on GaAs is not so remarkable as that on InP. It can be easily seen that many TDs such as 60° dislocations and 60° dislocation pairs generated in the interface of InAlAs MBL and GaAs substrate of sample S have penetrated through the InAlAs buffer and came into the active region. Roughly calculated from **Figure 19(a)**, the TDD in the absorption layer is more than 10<sup>9</sup> cm−2 in sample S, much higher than that of sample P with the TDD estimated to be less than 108 cm−2.

**Figure 19.** XTEM images of (a) sample S, (b) sample P, and (c) sample S0. Reprinted with permission from Elsevier.

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch

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223

3.6 3.7 3.8 3.9 4.0 4.1 4.2 4.3 4.4

1000/T(K)

EC-0.305 eV

GaAs-based PD

100 150 200 250 300 350

Temperature (K)

**Figure 20.** DLTS spectra for the two samples measured at a rate window of 680 Hz. Inset: Arrhenius plot of the deep state observed in GaAs-based PD sample using different lock-in frequencies of 41 Hz, 240 Hz, 320 Hz and 680 Hz. Reprinted

InP-based PD

GaAs-based PD

Electron Trap




ln(en/T2) (s-1

K-2

)




DLTS signal (arb. units)

with permission from Elsevier.



0.0

0.5

Second, deep level transient spectroscopy (DLTS) was measured in the temperature range of 77–300 K using DLS-83D DLTS test system and analyzed using standard techniques for the two samples. **Figure 20** shows the DLTS temperature scan signals using reverse bias voltage V0 = −2.0 V, filling pulse height VP = 0.5 V, and filling pulse duration tp = 20 μs. It is evident that no clear peak can be found in the scan signal of InP-based PD. This indicates a low enough density of electrically active defect in the active region of sample P. By contrast, a large electron trap peak around 275 K is clearly observed in that of sample S. Temperature scans with other three lock-in frequencies have also been made and put together to acquire emission


**Table 5.** The measured material and device results of samples A and B at 300 K (reprinted with permission from Elsevier).

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch http://dx.doi.org/10.5772/intechopen.70259 223

layer of sample S is much larger than that of sample P with values of 1638 and 644 arcsec, respectively. These suggests that the crystal quality of the In0.83Ga0.17As grown on the GaAs substrate is not so good as that grown on the InP substrate with a higher lattice mismatch by using the same type of MBL. This verified one of the conclusions in our previous work [31] that full relaxation and favorable optical property of the InGaAs layer could not occur for the PD wafers with a mismatch grading rate of 2.4% μm−1 or above but occur for the wafers with

Correspondingly, quite different temperature dependent I-V characteristics of the two PD chips with diameter of 300 μm were observed in the temperature range from 77 to 300 K, as shown in **Figure 18(b)**. At reverse bias of −10 mV, the dark currents of 2.28 μA at 300 K and 2.17 nA at 77 K for sample S are much larger than that of 674 nA at 300 K and 3.99 pA at 77 K

that of the InP-based PD at 300 K [32], as shown in **Table 5**. This indicates that via appropriate structural design, GaAs may act as a feasible substrate for replacement of InP for In0.83Ga0.17As

To analyze the cause of the different electrical performances of the two samples, XTEM measurement was performed firstly, as shown in **Figure 19**. From the XTEM bright-field images, it is obvious that the majority of misfit dislocations of sample P are mainly localized at the

vented from propagating into the In0.83Ga0.17As absorption layer by the continuously graded InAlAs MBL on InP under this strain gradient condition. However, it seems that the effect of the continuously graded InAlAs MBL on GaAs is not so remarkable as that on InP. It can be easily seen that many TDs such as 60° dislocations and 60° dislocation pairs generated in the interface of InAlAs MBL and GaAs substrate of sample S have penetrated through the InAlAs buffer and came into the active region. Roughly calculated from **Figure 19(a)**, the TDD in the

Second, deep level transient spectroscopy (DLTS) was measured in the temperature range of 77–300 K using DLS-83D DLTS test system and analyzed using standard techniques for the two samples. **Figure 20** shows the DLTS temperature scan signals using reverse bias voltage

 = −2.0 V, filling pulse height VP = 0.5 V, and filling pulse duration tp = 20 μs. It is evident that no clear peak can be found in the scan signal of InP-based PD. This indicates a low enough density of electrically active defect in the active region of sample P. By contrast, a large electron trap peak around 275 K is clearly observed in that of sample S. Temperature scans with other three lock-in frequencies have also been made and put together to acquire emission

> **R@1310 nm (A/W)**

**Table 5.** The measured material and device results of samples A and B at 300 K (reprinted with permission from Elsevier).

S 97.5 1.5 0.48 45 3.02 1.30 × 1010 P 97.7 0.5 0.56 53 9.07 2.25 × 1010

**ηe**

cm−2.

A of the GaAs-based PD is comparable with

Al1−xAs MBL. Threading dislocations have been pre-

cm−2 in sample S, much higher than that of sample P with

**@1310 nm (%) R0**

**A (Ω cm2 )** **D\* λp (cm Hz1/2/W)**

a mismatch grading rate of about 1.2% μm−1 or lower.

for sample P, respectively, while the zero bias R0

PDs in some low-end application area around RT.

early stage of the continuously graded Inx

absorption layer is more than 10<sup>9</sup>

V0

222 Epitaxy

the TDD estimated to be less than 108

**Sample Relax degree (%) Residual strain ε** 

**(10−3)**

**Figure 19.** XTEM images of (a) sample S, (b) sample P, and (c) sample S0. Reprinted with permission from Elsevier.

**Figure 20.** DLTS spectra for the two samples measured at a rate window of 680 Hz. Inset: Arrhenius plot of the deep state observed in GaAs-based PD sample using different lock-in frequencies of 41 Hz, 240 Hz, 320 Hz and 680 Hz. Reprinted with permission from Elsevier.

rate-temperature (en-T) data pairs. Then, the standard least-squares fitting for the plot (Arrhenius plot) was applied to extract the trap activation energy and the capture cross-section from the signal peaks. The results are shown in the inset of **Figure 20**. The trap activation energy E<sup>c</sup> − ET = 0.305 eV and the capture cross-section σn = 2.25 × 10−18 cm−2. This trap will definitely have a deleterious effect on the electrical performance of the PD. Previous literature has reported some deep centers located above the maximum of valence band in lattice-matched In0.53Ga0.47As/InP PDs, which were believed to be related with Fe impurities diffused from the Fe-doped substrates to the InGaAs layers [33]. Obviously, the trap occurred here cannot be resulted from Fe impurities because no Fe-doped substrate has been used in this experiment. However, similar traps have been observed in lattice-mismatched In1−xGax As/GaAs hetero-structures [34] and In0.78Ga0.22As/InP PDs [35]. Therefore, this 275 K electron trap signal could be associated with the dislocations and point defects in the In0.83Ga0.17As layer because of the even higher lattice mismatch on GaAs compared to that on InP. Since the trap level observed here locates near the center of the bandgap, it is considered that the trap-assisted tunneling current could be the main source of the large dark current in the GaAs-based PD at low temperature range.

It indicates that the dark current is dominated by the diffusion current (Idiff) with temperature

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch

perature range of 175–225 K, that is, the thermal activation energy is about half of the band gap energy of In0.83Ga0.17As. It means that the dark current is dominated by the generationrecombination current (Ig-r) in this range. As the temperature reduces lower than 175 K, E<sup>a</sup> becomes smaller, and thus, the components of trap-assisted tunneling current (Itat) and/or band-to-band tunneling (Ibtb) start to play an important role. However, it was found that the

PD. That is, Idiff and Ig-r dominate the dark current of the GaAs-based PD in a much smaller temperature region than that of InP-based PD. Alternatively, the majority component is tun-

In0.83Ga0.17As. This explained the higher dark current of the GaAs-based PD at the low-temperature region shown in **Figure 18(b)**. The large component proportion of tunneling current could be due to a high TDD in the active region caused by the relatively high lattice mismatch

To further optimize the GaAs-based PD structure and understand the strain relaxation mechanism with such a relatively high lattice mismatch, it is necessary to increase the buffer

Therefore, as a comparison, an In0.83Ga0.17As PD structure (renamed as sample S0) with a

layer has been also grown on a GaAs substrate. So that the mismatch grading rate in the

InP substrate. By comparison with samples S and P, the XRD and PL results of sample S0 are really better than those of sample S but still worse than that of sample P. Roughly estimated from the XTEM image of **Figure 19(c)**, the average TDD in the absorption layer of sample S0 is only slightly less than that of sample S but still much larger than that of sample P. In addition, some differences were observed between the two XTEM images of samples S and S0. The 60° dislocations and 60° dislocation pairs, which are dominant in sample S, have been mostly replaced by Lomer (90°) dislocations in sample S0. It would be resulted from reactions of 60° dislocations from different glide systems to form Lomer dislocations on the top of the InAlAs MBL due to a two-dimensional growth mode in the thicker buffer of sample S0 [36, 37]. This means that the lower mismatch grading rate of the MBL has improved the characteristics of

Therefore, it is supposed that the composition continuously graded buffer may not be a good strategy for growth of material systems with relatively high lattice mismatch due to a quite low growth rate of the MBE technique. A more suitable thin buffer layer should be exploited,

**3.2. Optimization of InAlAs metamorphic buffer on GaAs with relatively high lattice** 

From the discussions above, GaAs-based high In content InGaAs PDs may be not a good choice for practical special application of remote sensing at the low temperature range.

between In0.83Ga0.17As and GaAs [30]. This agrees well with the DLTS results.

thickness and thus decrease the strain gradient of the continuously graded Inx

Al1−xAs MBL was lowered to about 1.2% μm−1, close to that used for the Inx

the PD structure very limited except for the evolution of the TDs.

and the metamorphic strategy will be updated accordingly.

in the GaAs-based PD is much stronger than that of InP-based

Al1−xAs MBL and 1.5-μm-thick In0.83Ga0.17As absorption

is about 0.24 eV in the tem-

225

http://dx.doi.org/10.5772/intechopen.70259

is much less than the band gap of

Al1−xAs MBL.

Al1−xAs MBL on

ranging from about 225 to 300 K. As the temperature drops, E<sup>a</sup>

neling current at temperatures lower than 250 K, when E<sup>a</sup>

temperature dependence of Ea

5-μm-thick continuously graded In<sup>x</sup>

Inx

**mismatch**

On the other hand, the dark current could be expressed by using the thermal activation energy Ea and temperature T as Id ~ exp(−E<sup>a</sup> /kT) at a fixed reverse bias voltage, where E<sup>a</sup> = Eg /n, E<sup>g</sup> , and n are the band gap energy and the ideal factor, respectively. Therefore, to see the temperature-dependent characteristics of the dark current, Arrhenius plots of the dark currents at −10 mV are made for both samples, as shown in **Figure 21**. For the InP-based PD, near room temperature Ea is almost equivalent to the band gap energy of In0.83Ga0.17As (Eg = 0.48 eV).

**Figure 21.** Arrhenius plots of the dark current at −10 mV versus reciprocal temperatures of GaAs- and InP-based PDs. The activation energies in specific temperature regions are also shown in the figure. Reprinted with permission from Elsevier.

It indicates that the dark current is dominated by the diffusion current (Idiff) with temperature ranging from about 225 to 300 K. As the temperature drops, E<sup>a</sup> is about 0.24 eV in the temperature range of 175–225 K, that is, the thermal activation energy is about half of the band gap energy of In0.83Ga0.17As. It means that the dark current is dominated by the generationrecombination current (Ig-r) in this range. As the temperature reduces lower than 175 K, E<sup>a</sup> becomes smaller, and thus, the components of trap-assisted tunneling current (Itat) and/or band-to-band tunneling (Ibtb) start to play an important role. However, it was found that the temperature dependence of Ea in the GaAs-based PD is much stronger than that of InP-based PD. That is, Idiff and Ig-r dominate the dark current of the GaAs-based PD in a much smaller temperature region than that of InP-based PD. Alternatively, the majority component is tunneling current at temperatures lower than 250 K, when E<sup>a</sup> is much less than the band gap of In0.83Ga0.17As. This explained the higher dark current of the GaAs-based PD at the low-temperature region shown in **Figure 18(b)**. The large component proportion of tunneling current could be due to a high TDD in the active region caused by the relatively high lattice mismatch between In0.83Ga0.17As and GaAs [30]. This agrees well with the DLTS results.

rate-temperature (en-T) data pairs. Then, the standard least-squares fitting for the plot (Arrhenius plot) was applied to extract the trap activation energy and the capture cross-section from the signal peaks. The results are shown in the inset of **Figure 20**. The trap activation energy E<sup>c</sup>

0.305 eV and the capture cross-section σn = 2.25 × 10−18 cm−2. This trap will definitely have a deleterious effect on the electrical performance of the PD. Previous literature has reported some deep centers located above the maximum of valence band in lattice-matched In0.53Ga0.47As/InP PDs, which were believed to be related with Fe impurities diffused from the Fe-doped substrates to the InGaAs layers [33]. Obviously, the trap occurred here cannot be resulted from Fe impurities because no Fe-doped substrate has been used in this experiment. However, similar

In0.78Ga0.22As/InP PDs [35]. Therefore, this 275 K electron trap signal could be associated with the dislocations and point defects in the In0.83Ga0.17As layer because of the even higher lattice mismatch on GaAs compared to that on InP. Since the trap level observed here locates near the center of the bandgap, it is considered that the trap-assisted tunneling current could be the

On the other hand, the dark current could be expressed by using the thermal activation energy

and n are the band gap energy and the ideal factor, respectively. Therefore, to see the temperature-dependent characteristics of the dark current, Arrhenius plots of the dark currents at −10 mV are made for both samples, as shown in **Figure 21**. For the InP-based PD, near room

is almost equivalent to the band gap energy of In0.83Ga0.17As (Eg

Temperature (K)

4 6 8 10 12

**Figure 21.** Arrhenius plots of the dark current at −10 mV versus reciprocal temperatures of GaAs- and InP-based PDs. The activation energies in specific temperature regions are also shown in the figure. Reprinted with permission from

1000/T (K-1

)

=0.09 eV

Ea

Ea

=0.48 eV

=0.21 eV

Ea

Ea

Ea

Ea

=0.48 eV

=0.10 eV

300 250 200 150 100

=0.24 eV

/kT) at a fixed reverse bias voltage, where E<sup>a</sup>

main source of the large dark current in the GaAs-based PD at low temperature range.

traps have been observed in lattice-mismatched In1−xGax

and temperature T as Id ~ exp(−E<sup>a</sup>

1E-11

1E-10

I

Elsevier.

D (A) @ -10 mV

1E-9

1E-8

1E-7

1E-6

Ea

224 Epitaxy

temperature Ea

− ET =

As/GaAs hetero-structures [34] and

InP-based PD

GaAs-based PD

= Eg

/n, E<sup>g</sup> ,

= 0.48 eV).

To further optimize the GaAs-based PD structure and understand the strain relaxation mechanism with such a relatively high lattice mismatch, it is necessary to increase the buffer thickness and thus decrease the strain gradient of the continuously graded Inx Al1−xAs MBL. Therefore, as a comparison, an In0.83Ga0.17As PD structure (renamed as sample S0) with a 5-μm-thick continuously graded In<sup>x</sup> Al1−xAs MBL and 1.5-μm-thick In0.83Ga0.17As absorption layer has been also grown on a GaAs substrate. So that the mismatch grading rate in the Inx Al1−xAs MBL was lowered to about 1.2% μm−1, close to that used for the Inx Al1−xAs MBL on InP substrate. By comparison with samples S and P, the XRD and PL results of sample S0 are really better than those of sample S but still worse than that of sample P. Roughly estimated from the XTEM image of **Figure 19(c)**, the average TDD in the absorption layer of sample S0 is only slightly less than that of sample S but still much larger than that of sample P. In addition, some differences were observed between the two XTEM images of samples S and S0. The 60° dislocations and 60° dislocation pairs, which are dominant in sample S, have been mostly replaced by Lomer (90°) dislocations in sample S0. It would be resulted from reactions of 60° dislocations from different glide systems to form Lomer dislocations on the top of the InAlAs MBL due to a two-dimensional growth mode in the thicker buffer of sample S0 [36, 37]. This means that the lower mismatch grading rate of the MBL has improved the characteristics of the PD structure very limited except for the evolution of the TDs.

Therefore, it is supposed that the composition continuously graded buffer may not be a good strategy for growth of material systems with relatively high lattice mismatch due to a quite low growth rate of the MBE technique. A more suitable thin buffer layer should be exploited, and the metamorphic strategy will be updated accordingly.

#### **3.2. Optimization of InAlAs metamorphic buffer on GaAs with relatively high lattice mismatch**

From the discussions above, GaAs-based high In content InGaAs PDs may be not a good choice for practical special application of remote sensing at the low temperature range. However, it is still valuable to design and develop an appropriate buffer scheme for the device development from material system with relatively high-lattice mismatch.

Since that the way strain is introduced at the initial stage of the MBL has been proved to play a critical role in the final TDD [10]. The strain energy should be released as quickly as possible, and the multiplication of TD must be avoided occurring at the final stage of the buffer layer. Therefore, if we take the accessory advantage of this relatively high lattice mismatch between high In content InGaAs and GaAs, we can promote the nucleation of dispersed TDs at the initial stage of the MBL. Considering that the high In content InAlAs and/or InGaAs will quickly achieve a high relaxation on GaAs due to the small critical thickness. Specifically, for the growth of In0.83Ga0.17As PD on GaAs, we can use the fixed-composition In0.83Al0.17As as MBL to accelerate the release of strain energy and the nucleation of dispersed misfit dislocations at the initial stage of the buffer and thus restrain the misfit dislocation at the interface between In0.83Al0.17As and GaAs. The uniform composition buffer will also reduce the strain gradient at the later stage of the MBL with lattice nearly fully relaxed. We show that this method results in a lower TDD and a smoother surface of the final absorption and cap layer.

In this work, four In0.83Ga0.17As PD structures with different buffer schemes were grown by GSMBE on S.I. (1 0 0)-oriented GaAs epi-ready substrates [38]. Each structure consisted of a 2.5-μm N+ InAlAs MBL and a 1.5-μm n− In0.83Ga0.17As absorption layer followed by a 530 nm P+ InAlAs cap. The detailed buffer schemes for samples GS1, GS2, GS3, and GS4 were listed in **Table 6**. Growth condition was exactly the same as that in our previous study [30]. The strategies of substrate temperature graded from 530 to 460°C, and the uniform temperature of 490 °C were adopted for the deposition of continuously graded Inx Al1−xAs and fixed-composition In0.83Al0.17As buffer, respectively. In addition, InAs wetting layer was inserted between the fixed-composition In0.83Al0.17As layer and the GaAs substrate for samples GS3 and GS4 to investigate the effect of interfacial layer on the TD behaviors at the interface.

As shown in **Figure 22**, compared to sample GS1, the cross-hatch pattern in the 2D AFM image of sample GS2 is less pronounced, and smaller diameters of 3D mounds align along the [1 1 0] direction in the 3D AFM image, indicating a smaller residual strain fields on the surface [39]. Thus, the RMS roughnesses of samples GS2–GS4 are a little smaller than that of sample GS1, as summarized in **Table 6**. As well known, when the thickness of the epilayer is beyond the critical value, the misfit strain could relax by introducing misfit dislocation arrays as well as surface undulation [40, 41]. The larger RMS value and larger size of 3D mounds on the structural surface of sample GS1 reflect that the strain-reliving surface roughening is more serious in sample GS1 than samples GS2–GS4.

**Figure 23(a)** shows the normalized HRXRD rocking curves for (0 0 4) reflections of all samples, which is also related to the TDD and other crystal imperfections [42]. A broader peak located at about 31° of sample GS1 reflects the continuously grading profile of In<sup>x</sup> Al1−xAs (x = 0.1 → 0.86) MBL which ultimately combined with the In0.83Ga0.17As peak. While much lower FWHM of In0.83Ga0.17As layers for samples GS2–GS4 mainly due to the thicker fixedcomposition In0.83Al0.17As MBL. Though all PD structures exhibited high degree of lattice relaxation, shown in **Table 6**, the improvement of structural surface and crystal quality was believed to be associated with the amelioration of misfit dislocation formation at the

**Figure 22.** AFM images of samples A and B. The scan area is 20 × 20 μm<sup>2</sup>

**Sample Buffer scheme XRD FWHM** 

grades from 0.1 to 0.86

GS3 5 MLs InAs QDs + 2.5 μm In0.83Al0.17As

GS4 50 nm InAs + 2.5 μm In0.83Al0.17As

Al1−xAs buffer, x

GS1 2.5 μm In<sup>x</sup>

**(arcsec)**

GS2 2.5 μm In0.83Al0.17As 544 97.5 0.82 8.1

**Table 6.** Buffer schemes and measured results for four samples (reprinted with permission from Elsevier).

**Relaxation degree** 

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch

1782 94.8 0.84 11.6

410 98.3 0.83 8.8

511 98.8 0.82 8.6

**In composition AFM RMS** 

http://dx.doi.org/10.5772/intechopen.70259

**(nm)**

227

**(%)**

. Reprinted with permission from Elsevier.

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch http://dx.doi.org/10.5772/intechopen.70259 227


**Table 6.** Buffer schemes and measured results for four samples (reprinted with permission from Elsevier).

However, it is still valuable to design and develop an appropriate buffer scheme for the device

Since that the way strain is introduced at the initial stage of the MBL has been proved to play a critical role in the final TDD [10]. The strain energy should be released as quickly as possible, and the multiplication of TD must be avoided occurring at the final stage of the buffer layer. Therefore, if we take the accessory advantage of this relatively high lattice mismatch between high In content InGaAs and GaAs, we can promote the nucleation of dispersed TDs at the initial stage of the MBL. Considering that the high In content InAlAs and/or InGaAs will quickly achieve a high relaxation on GaAs due to the small critical thickness. Specifically, for the growth of In0.83Ga0.17As PD on GaAs, we can use the fixed-composition In0.83Al0.17As as MBL to accelerate the release of strain energy and the nucleation of dispersed misfit dislocations at the initial stage of the buffer and thus restrain the misfit dislocation at the interface between In0.83Al0.17As and GaAs. The uniform composition buffer will also reduce the strain gradient at the later stage of the MBL with lattice nearly fully relaxed. We show that this method results

In this work, four In0.83Ga0.17As PD structures with different buffer schemes were grown by GSMBE on S.I. (1 0 0)-oriented GaAs epi-ready substrates [38]. Each structure consisted of a

 InAlAs cap. The detailed buffer schemes for samples GS1, GS2, GS3, and GS4 were listed in **Table 6**. Growth condition was exactly the same as that in our previous study [30]. The strategies of substrate temperature graded from 530 to 460°C, and the uniform temperature of

sition In0.83Al0.17As buffer, respectively. In addition, InAs wetting layer was inserted between the fixed-composition In0.83Al0.17As layer and the GaAs substrate for samples GS3 and GS4 to

As shown in **Figure 22**, compared to sample GS1, the cross-hatch pattern in the 2D AFM image of sample GS2 is less pronounced, and smaller diameters of 3D mounds align along the [1 1 0] direction in the 3D AFM image, indicating a smaller residual strain fields on the surface [39]. Thus, the RMS roughnesses of samples GS2–GS4 are a little smaller than that of sample GS1, as summarized in **Table 6**. As well known, when the thickness of the epilayer is beyond the critical value, the misfit strain could relax by introducing misfit dislocation arrays as well as surface undulation [40, 41]. The larger RMS value and larger size of 3D mounds on the structural surface of sample GS1 reflect that the strain-reliving surface roughening is more

**Figure 23(a)** shows the normalized HRXRD rocking curves for (0 0 4) reflections of all samples, which is also related to the TDD and other crystal imperfections [42]. A broader peak

(x = 0.1 → 0.86) MBL which ultimately combined with the In0.83Ga0.17As peak. While much lower FWHM of In0.83Ga0.17As layers for samples GS2–GS4 mainly due to the thicker fixedcomposition In0.83Al0.17As MBL. Though all PD structures exhibited high degree of lattice relaxation, shown in **Table 6**, the improvement of structural surface and crystal quality was believed to be associated with the amelioration of misfit dislocation formation at the

located at about 31° of sample GS1 reflects the continuously grading profile of In<sup>x</sup>

In0.83Ga0.17As absorption layer followed by a 530 nm

Al1−xAs and fixed-compo-

Al1−xAs

development from material system with relatively high-lattice mismatch.

in a lower TDD and a smoother surface of the final absorption and cap layer.

490 °C were adopted for the deposition of continuously graded Inx

investigate the effect of interfacial layer on the TD behaviors at the interface.

InAlAs MBL and a 1.5-μm n−

serious in sample GS1 than samples GS2–GS4.

2.5-μm N+

P+

226 Epitaxy

**Figure 22.** AFM images of samples A and B. The scan area is 20 × 20 μm<sup>2</sup> . Reprinted with permission from Elsevier.

**Figure 23.** (a) (0 0 4) *ω* − 2*θ* high resolution XRD patterns. (b) PL spectra at 300 K and 10 K of samples GS1–GS4. Reprinted with permission from Elsevier.

layer because of the relatively high-lattice mismatch of 5.9% with respect to GaAs. Therefore, the relaxation process in the fixed-composition buffer of sample GS2 may start with a layerto-island transition (Stranski-Kranstanov growth mode) after the deposition of a couple MLs of In0.83Al0.17As on GaAs. Then, the large compressive strain can be released by generation and reaction of misfit dislocation networks. Since the initial growth mechanism will also affect the

Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch

http://dx.doi.org/10.5772/intechopen.70259

229

**Figure 24.** XTEM images of (a) sample GS1 and (b) sample GS2. Reprinted with permission from Elsevier.

lattice mismatch [10]. The formation of In0.83Al0.17As QDs at the initial stage of the In0.83Al0.17As MBL has indeed played an important role by the formation of high density of nucleation sites [13], which can act as core of misfit dislocations. In this way, the relaxation process is promptly due to the substantial lattice mismatch and most of the misfit dislocations are inhibited to the core areas, which locate close to the In0.83Al0.17As/GaAs interface. Therefore, TDs in the buffer of sample GS2 cannot propagate a long distance due to the island growth near the interface. They annihilated much easier by interaction with the point defects existing in

expected that the intermediate layer of 5 MLs InAs QD will enhance the dislocation nucleation process at the In0.83Al0.17As/GaAs interface because of a higher lattice mismatch between InAs and GaAs. Through appropriate buffer design, fixed-composition In0.83Al0.17As buffer may be a better choice for the growth of metamorphic In0.83Ga0.17As PDs on GaAs substrate.

In conclusion, due to the inventions of metamorphic techniques such as, linearly, step, and one-step continuously graded buffer, compositional overshoot, and DA in the growth of

Ga1−xAs on GaAs except for the

Al1−xAs MBL of sample GS1. It is

dislocation generation mechanism in the hetero-epitaxy of Inx

the high In content In0.83Al0.17As layer with respect to the Inx

**4. Conclusion**

film/substrate interface [43]. This was further demonstrated by sample GS3 with a narrowest FWHM value of 410 arcs by inserting 5 MLs InAs QDs at the In0.83Al0.17As/GaAs interface. This indicates that the growth of the QD is beneficial to the dislocation nucleation at the interface and the confinement of misfit dislocations in the inactive region, while the effect is slightly weakened when InAs was thickened and degenerated, resulting in a degraded interfacial quality.

From the PL spectra measured at 300 and 10 K shown in **Figure 23(b)**, we can see that the PL intensity of the In0.83Ga0.17As absorption layers increased in sequence for samples GS1–GS3, meaning that the crystal defect has been suppressed through the substitution of continuously graded Inx Al1−xAs MBL by fixed-composition In0.83Al0.17As MBL. The XTEM images of the two typical PD structures were shown in **Figure 24**. For sample GS1, it is notable that the misfit dislocation networks were separated by 0.1 μm from the interface between In0.1Al0.9As and GaAs because of the In0.1Al0.9As layer, and highly intensive dislocation arrays accumulated in the lower area of the Inx Al1−xAs grading layer, as shown in **Figure 24(a)**. Many vertical dislocations with long dislocation length, thread from the tangling area through the buffer along the [1 0 0] direction and propagate to the top active region of the PD structure. Strain in this structure released dominantly by surface undulation, which deteriorates the surface morphology of the PD structure and leads to a large RMS roughness. The surface ripple troughs–induced misfit dislocations were generated by gliding of dislocation half loops from the surface through the epilayer [44].

However, the dislocations in the structure of sample GS2 are short, and most of them are not perpendicular to the sample surface, as shown in **Figure 24(b)**. We supposed that In0.83Al0.17As QDs would form at the beginning of growth process of the fixed-composition In0.83Al0.17As Epitaxy and Device Properties of InGaAs Photodetectors with Relatively High Lattice Mismatch http://dx.doi.org/10.5772/intechopen.70259 229

**Figure 24.** XTEM images of (a) sample GS1 and (b) sample GS2. Reprinted with permission from Elsevier.

layer because of the relatively high-lattice mismatch of 5.9% with respect to GaAs. Therefore, the relaxation process in the fixed-composition buffer of sample GS2 may start with a layerto-island transition (Stranski-Kranstanov growth mode) after the deposition of a couple MLs of In0.83Al0.17As on GaAs. Then, the large compressive strain can be released by generation and reaction of misfit dislocation networks. Since the initial growth mechanism will also affect the dislocation generation mechanism in the hetero-epitaxy of Inx Ga1−xAs on GaAs except for the lattice mismatch [10]. The formation of In0.83Al0.17As QDs at the initial stage of the In0.83Al0.17As MBL has indeed played an important role by the formation of high density of nucleation sites [13], which can act as core of misfit dislocations. In this way, the relaxation process is promptly due to the substantial lattice mismatch and most of the misfit dislocations are inhibited to the core areas, which locate close to the In0.83Al0.17As/GaAs interface. Therefore, TDs in the buffer of sample GS2 cannot propagate a long distance due to the island growth near the interface. They annihilated much easier by interaction with the point defects existing in the high In content In0.83Al0.17As layer with respect to the Inx Al1−xAs MBL of sample GS1. It is expected that the intermediate layer of 5 MLs InAs QD will enhance the dislocation nucleation process at the In0.83Al0.17As/GaAs interface because of a higher lattice mismatch between InAs and GaAs. Through appropriate buffer design, fixed-composition In0.83Al0.17As buffer may be a better choice for the growth of metamorphic In0.83Ga0.17As PDs on GaAs substrate.

#### **4. Conclusion**

film/substrate interface [43]. This was further demonstrated by sample GS3 with a narrowest FWHM value of 410 arcs by inserting 5 MLs InAs QDs at the In0.83Al0.17As/GaAs interface. This indicates that the growth of the QD is beneficial to the dislocation nucleation at the interface and the confinement of misfit dislocations in the inactive region, while the effect is slightly weakened when InAs was thickened and degenerated, resulting in a degraded interfacial quality.

**Figure 23.** (a) (0 0 4) *ω* − 2*θ* high resolution XRD patterns. (b) PL spectra at 300 K and 10 K of samples GS1–GS4.

From the PL spectra measured at 300 and 10 K shown in **Figure 23(b)**, we can see that the PL intensity of the In0.83Ga0.17As absorption layers increased in sequence for samples GS1–GS3, meaning that the crystal defect has been suppressed through the substitution of continuously

typical PD structures were shown in **Figure 24**. For sample GS1, it is notable that the misfit dislocation networks were separated by 0.1 μm from the interface between In0.1Al0.9As and GaAs because of the In0.1Al0.9As layer, and highly intensive dislocation arrays accumulated in the lower

long dislocation length, thread from the tangling area through the buffer along the [1 0 0] direction and propagate to the top active region of the PD structure. Strain in this structure released dominantly by surface undulation, which deteriorates the surface morphology of the PD structure and leads to a large RMS roughness. The surface ripple troughs–induced misfit dislocations were generated by gliding of dislocation half loops from the surface through the epilayer [44].

However, the dislocations in the structure of sample GS2 are short, and most of them are not perpendicular to the sample surface, as shown in **Figure 24(b)**. We supposed that In0.83Al0.17As QDs would form at the beginning of growth process of the fixed-composition In0.83Al0.17As

Al1−xAs MBL by fixed-composition In0.83Al0.17As MBL. The XTEM images of the two

Al1−xAs grading layer, as shown in **Figure 24(a)**. Many vertical dislocations with

graded Inx

228 Epitaxy

Reprinted with permission from Elsevier.

area of the Inx

In conclusion, due to the inventions of metamorphic techniques such as, linearly, step, and one-step continuously graded buffer, compositional overshoot, and DA in the growth of lattice-mismatched material system, the development of InP- and GaAs-based wavelengthextended Inx Ga1−xAs (x > 0.53) photodetectors with relatively high lattice mismatch has achieved remarkable success in SWIR band especially for the applications at higher operation temperatures and robust circumstances. Though GaAs-based In0.83Ga0.17As PD shows large potential in some low-end application area around RT, the high densities of TDs and electron traps in the active region due to the relatively high-lattice mismatch have hindered its development. By designing of abrupt interface and appropriate buffer, the initial dislocation nucleation process may pave a way of development of lattice-mismatched device structures.

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## **Acknowledgements**

The authors wish to acknowledge the support of the National Key Research and Development Program of China under Grant No. 2016YFB0402400, the National Natural Science Foundation of China under grant Nos. 61405232, 61675225, and 61605232, and the Youth Innovation Promotion Association CAS under Grant No. 2013155.

## **Author details**

Xing-you Chen\*, Yi Gu and Yong-gang Zhang

\*Address all correspondence to: xychen@mail.sim.ac.cn

State Key Laboratory of Functional Materials for Informatics, Shanghai Institute of Microsystem and Information Technology, Chinese Academy of Sciences, China

## **References**


[5] Ji L, Lu SL, Zhao YM, Tan M, Dong JR, Yang H. Mint: Compositionally undulating step-graded InAsy P1−y buffer layer growth by metal-organic chemical vapor deposition. Journal of Crystal Growth. 2013;**363**:44-48. DOI: 10.1016/j.jcrysgro.2012.09.035

lattice-mismatched material system, the development of InP- and GaAs-based wavelength-

achieved remarkable success in SWIR band especially for the applications at higher operation temperatures and robust circumstances. Though GaAs-based In0.83Ga0.17As PD shows large potential in some low-end application area around RT, the high densities of TDs and electron traps in the active region due to the relatively high-lattice mismatch have hindered its development. By designing of abrupt interface and appropriate buffer, the initial dislocation nucleation process may pave a way of development of lattice-mismatched device structures.

The authors wish to acknowledge the support of the National Key Research and Development Program of China under Grant No. 2016YFB0402400, the National Natural Science Foundation of China under grant Nos. 61405232, 61675225, and 61605232, and the Youth Innovation

State Key Laboratory of Functional Materials for Informatics, Shanghai Institute of Microsystem and Information Technology, Chinese Academy of Sciences, China

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Ga1−xAs (x > 0.53) photodetectors with relatively high lattice mismatch has

extended Inx

230 Epitaxy

**Acknowledgements**

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## *Edited by Miao Zhong*

The edited volume "Epitaxy" is a collection of reviewed and relevant research chapters, offering a comprehensive overview of recent developments in the field of materials science. The book comprises single chapters authored by various researchers and edited by an expert active in this research area. All chapters are complete in themselves but are united under a common research study topic. This publication aims at providing a thorough overview of the latest research efforts by international authors in the field of materials science as well as opening new possible research paths for further developments.

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Epitaxy

Epitaxy