Hexagonal Nanopyramidal Prisms of Nearly Intrinsic InN on Patterned

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Hexagonal nano-pyramidal-prisms of nearly intrinsic InN on patterned GaN nanowire arrays A.T.M. Golam Sarwar, Benjamin Leung, George T Wang, and Roberto C Myers Cryst. Growth Des., Just Accepted Manuscript • DOI: 10.1021/acs.cgd.7b01725 • Publication Date (Web): 04 Jan 2018 Downloaded from http://pubs.acs.org on January 5, 2018

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Crystal Growth & Design

Hexagonal nano-pyramidal-prisms of nearly intrinsic InN on patterned GaN nanowire arrays ATM Golam Sarwar1, Benjamin Leung2, George T Wang2, and Roberto C Myers1,3* 1

Department of Electrical and Computer Engineering, The Ohio State University, Columbus, Ohio 43210, USA. 2

3

Sandia National Laboratories, Albuquerque, New Mexico, 87185, USA

Department of Materials Science and Engineering, The Ohio State University, Columbus, Ohio 43210, USA.

ABSTRACT: By using multiple growth steps that separate the nucleation and growth processes, we show that nearly intrinsic InN single nanocrystals of high optical quality can be formed on patterned GaN nanowire arrays by molecular beam epitaxy. The InN nanostructures form into well-defined hexagonal prisms with pyramidal tops. Micro-photoluminescence (µ-PL) is carried out at low temperature (LT: 28.2 K) and room temperature (RT: 285 K) to gauge the relative material quality of the InN nanostructures. Nano-pyramidal-prisms grown using a three-step growth method are found to show superior quantum efficiency. Excitation and temperature dependent µ-PL demonstrates the very high quality and nearly intrinsic nature of the ordered InN nanostructure arrays.

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The III-Nitride materials system has been an area of increasing interest due to their direct bandgaps spanning from the deep ultraviolet (DUV) to the infrared (IR). InN in particular has potential applications in chemical sensing, IR optoelectronics, and high efficiency solar cells [1] [2] [3]due to having a bandgap of 1770 nm) [4] and radiation resistance [5]. Among the III-Nitrides, it also has the smallest electron mass, resulting in a measured and predicted room temperature carrier mobilities greater than 1000 and 4400 cm2/Vs, respectively [6] [7]. However, its bandgap is highly susceptible to Moss-Burnstein shift [8] [9]This shift has proven difficult to remove, due to the difficulties in growing material free of charged defects [8] The growth of high quality InN thin films is hindered by high lattice mismatch (~11%) to GaN. Moreover, the growth kinetics of InN significantly change with temperature making it more difficult to control morphology and material quality than for GaN [10] [11] In an effort to circumvent these problems the use of nanostructures has been investigated, and InN nanowires have been demonstrated on Si, AlN, GaN, and even brass [12] [13] [14] [15] [16]. These provide pure, high quality material, but lack the order and convenience of thin films. Nanowires have also been investigated as strain-compliant templates for the coalescence of reduced defect density films [17] [18]. However, the wurtzite InN nanowires showed multiple crystalline rotations when grown on substrates with lower hexagonal symmetry, such as Si [13]. Here, we explore the use of a top-down fabricated GaN nanowire template as a potential substrate for the growth of high material-quality InN.. The patterned GaN nanowires are fabricated via a two-step, top–down dry plus wet etch process using c-plane (0001) oriented n-type GaN epilayers grown on sapphire by metal organic chemical vapor deposition (MOCVD) [19] [20]. A self-assembled mask of monodispersed silica


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spheres (typically between 500 nm and 3mm in diameter) is deposited [21] followed by an chlorine-based inductively coupled plasma (ICP) etching to create an array of tapered pillars. A KOH-based wet chemical etching (AZ400K photoresist developer) is performed to selectively etch the semipolar sidewalls to form nontapered nanowires with straight, vertical sidewalls and remove the dry-etched damaged sidewall material. Figure. 1 shows a scanning electron microscopy (SEM) image of GaN nanowires formed using this two-step top-down etch process. The average nanowire height is 1.6 µm with a 300 nm – 400 nm diameter and a 1 µm pitch.

Figure 1. Scanning electron micrograph of the GaN nanowires using two-step top-down etch process.

InN is subsequently grown on the GaN nanowires by plasma-assisted molecular beam epitaxy (PAMBE) in a Veeco GEN 930 system having an average base pressure of ~7×10-11 torr. The GaN nanowire templates are first cleaned in piranha solution (3 H2SO4:1 H2O2), followed by standard solvents before transfer to the MBE growth chamber. The InN is grown on the GaN nanowire templates under N-rich conditions. The In-flux is measured to be 7.8×10-8 torr using a


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beam flux monitor ion gauge. The N2 partial pressure and power is kept constant at 2×10-5 torr and 350 W, respectively. Table 1 summarizes all the samples grown in this study. (a) Sample A

(b) Sample B

InN NW on top surface

InN NW on side wals

Merging

Tsub = 450 °C

Tsub = 425 °C

Figure 2. Scanning electron micrographs of InN on patterned GaN nanowires at different growth conditions. Scale bars are 1µm.

Figure. 2(a) shows an SEM image of sample A grown at a substrate temperature (Tsub) of 425°C for 3 hours. Under these growth conditions, nucleation of multiple irregularly shaped InN nanowires (100 – 250 nm in diameter) proceeds on the top surface of each GaN nanowire. As they grow longer, these InN nanowires merge together and begin to form large faceted crystals. Large numbers of tilted small diameter nanowires are apparent on the sidewalls. By increasing Tsub to 450 °C (sample B), keeping all other parameters fixed, the InN nanowire density on the sidewalls is greatly reduced (Fig. 2(b)). Only single large InN hexagonal nano-prisms are observed on the top surfaces of the GaN nanowires (inset of Fig. 2(b)). The nano-prisms have larger diameters (300-700 nm) compared to the multiple sidewall InN nanowires observed in sample A. This is expected because the critical size of stable InN nuclei increases with


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temperature. In addition, the adatom mobility of In is increased leading to a higher effective In flux on the top c-plane of the GaN nanowires than on the sidewalls.

(a)

In

N

N

InN NW GaN NW Step 1

Step 2

1010

(b) Sample C

Figure 3. (a) Growth flow schematic and (b) Scanning electron micrograph of InN on patterned GaN nanowires using two-step growth process. Scales bars are 500 nm.

Next, we grow InN NWs employing a two-step growth process (sample C), wherein we first deposit InN on patterned GaN NWs using a Tsub of 425°C (same as sample A). In the second


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step, Tsub is increased to 460°C, keeping the N-plasma on, while closing the In shutter. Fig. 3(a) illustrates a simplified schematic of the growth process of sample C. Fig. 3(b) shows a tilted view SEM image of sample grown using the two-step method (sample C). A single InN nanowire (~350 – 450 nm diameter) is observed on the top surface of each GaN nanowire. It is worth noting, that although the first growth step for sample C is identical to that of sample A, the formation of multiple InN NWs on top of a single GaN nanowire was not observed. Therefore, it is clear that during the second step, the multiple small InN nanowires observed on sample A (Fig. 2(a)), merge together to form one large diameter InN nanowire via Ostwald ripening. Evidence of such ripening was previously reported in the self-assembled growth of InN on Si [13]. This merging process occurs because the small InN nanowires (on the top surface and sidewalls of GaN nanowires) are thermodynamically unstable at elevated temperature (Tsub = 460°C). Thus, these unstable InN nanowires decompose into In and N2. While decomposed N2 leaves the growth surface, In stays because of its high desorption temperature. The higher temperature (Tsub = 460 °C) allows for higher adatom mobility of the decomposed In, helping them migrate and incorporate at the growth front. For the same reason, the sidewalls of the GaN nanowires in sample C (Fig. 3(b)) also have fewer InN nanostructures compared to those in sample A (Fig. 2(a)). The inset of Fig. 3(b) shows a magnified image of an InN nanowire with an obvious hexagonal prism shape with only minor irregularities. The change in nanowire structural morphology in sample C led to the investigation of nanowire growth using a three-step growth method (sample D). Here, we first grow InN using


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the two steps identical to sample C followed by InN deposition for an additional 30 minutes at (a) Sample D

1 µm

(b)

Tsub = 500°C. (c) Figure 4. Plan view (a)-(b), and tilted view (c) scanning electron micrographs of InN on patterned GaN nanowires using three-step growth process.


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Nearly-perfect hexagonal InN nano-prisms are observed in the plan view SEM image of sample D (see Fig. 4(a)). A noticeable increase in diameter of nearly 2× (700 – 800 nm) is observed compared to sample C. Fig. 4(b) shows a magnified SEM image corresponding to the red dashed box of Fig. 4(a). It is clearly seen that the top surface of the InN nano-prisms are not flat, rather, there are six tilted planes forming a hexagonal pyramid top surface. Fig. 4(c) shows tilted view SEM image of sample D where this top hexagonal pyramid shape is more evident. In contrast to the previously discussed samples, sample D exhibits no apparent nanowires on the sidewalls. This is attributed to the enhanced decomposition, diffusion and incorporation of InN at the growth front at the even higher growth temperature during the third step. It is also important to realize that using the three-step growth method we are able to synthesis InN nano-pyramidalprisms at very high temperature. One step growth at such temperatures (Tsub = 500°C) usually gives a negligible growth rate of InN. This is because at such temperatures, nucleation becomes

101multi-step 0 unlikely. In the growth process, nucleation is performed at low temperatures. InN growth at high temperature then becomes possible due to the presence of pre-existing InN nuclei allowing growth while preventing additional nucleation. As the growth is performed closer to equilibrium, semi-polar faceting is observed due to the lower surface energy of those planes. As described below, these single crystal hexagonal InN nano-prisms exhibit remarkably high quality optical properties consistent with low residual electron concentration. Figure. 5 shows micro-photoluminescence (µ-PL) spectra of samples C and D at low temperature (LT) (28.2 K) and room temperature (RT) (285 K). We observe a substantial increase in the PL peak intensity (Ipeak) and a blue shift in the peak emission energy (Epeak) of sample D compared to that of sample C. Together this indicates that the material quality of


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sample D is better than that of sample C. We attribute this to the higher Tsub during the InN growth during the third step of sample D. During this step, the volume of InN nano-prisms increases, not only through deposition from the vapor phase, but also from In adatoms supplied by high temperature decomposition of InN. We also observe a red shift of Epeak for both samples with the increase of temperature from LT to RT.

12k Sample C Sample D

PL intensity (a. u.)

T = 28.2 K 10k

Excitation: 1.66 eV (750 nm) 21 mW

8k 6k 4k 2k 0 5k T = 285 K

PL intensity (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Excitation: 1.66 eV (750 nm) 25 mW

4k

Laser 2nd order 0.828 eV (1500 nm)

3k 2k 1k 0 0.55

0.60

0.65

0.70

0.75

0.80

0.85

Enegry (eV) Figure 5. Micro-photoluminescence spectra at 28.2 K (top) and 285 K (bottom) of InN on patterned GaN nanowire using different growth method.


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The superior material quality of sample D persuaded us to perform excitation and temperature dependent µ-PL on this sample. Figure 6(a) shows normalized PL spectra for different excitation power at ~28.2 K. The vertical shift is done for visual clarity. We measure Epeak of 0.683 eV and full width at half maximum (FWHM) of 26.5 meV at an excitation power of 10 µW. Epeak and FWHM (at this temperature) are in close agreement with reported values for self-assembled nearly intrinsic InN nanowires [22] [23] and much lower compared to n-type degenerate InN nanowires [24] [25]. This indicates the presence of small residual electron concentration in these InN nano-pyramidal-prisms grown on patterned GaN nanowires. We also observe an increase in Epeak (Fig.6(b)) and FWHM (Fig. 6(c)) with increasing excitation power. We observe that the Epeak increases slowly from low excitation power up to ~1.8 mW. At this point it starts to increase rapidly with the increase of excitation power. At ~20 mW it changes the slope again and starts to increase slowly. This creates an S-shape of the peak positions with increasing excitation power. Increasing the excitation power from 10 µW to 42.3 mW shows an increase in Epeak from 0.683 to 0.695 eV, and the FWHM increases from 26.5 to 44.6 meV. This happens due to the classic band filling effect in non-degenerate or intrinsic semiconductors [22] [23]. It is also worth noting that the increase in FWHM occurs mainly due to broadening of the high energy tail of the PL spectra (Fig. 6(a)) indicative of the increased electron temperature at higher pump powers. The inset of Fig. 6(b) shows the half intensity energy (the energy at which the PL intensity drops to half of Ipeak) at low (ELOW) and high (EHIGH) energy side. It is clearly evident that with the increase of excitation power(photo-excited carrier density) EHIGH increases slowly for small excitation power and starts to increase rapidly when excitation power is greater than ~1.8 mW. This is a direct observation of fermi energy (EF) shift into the


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conduction band with the increase of photo-excited carriers. This is also reflected in the shift of

T = 28.2 K 42.3 mW 30.2 mW 21.5 mW 12.0 mW 2.0 mW 0.6 mW 0.3 mW 0.2 mW 40 µW 10 µW

0.60

0.65

0.70

0.75

0.700 0.696 0.692

0.72

Energy (eV)

(a)

Peak Energy (eV)

Normalized PL (a.u.)

Epeak with increasing excitation power (Fig. 2(b)).

(c)

40 35 30 25 20

0.01

0.1

1

10

Power (mW)

(b)

EHIGH

0.68

0.66

0.688

0.01

0.1

1

10

Power (mW)

0.684

T = 28.2 K 0.01

0.80

0.1

1

10

Power (mW) Peak Intensity (a.u.)

45

ELOW

0.70

Energy (eV) 50

FWHM (meV)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


104

(d) ~P

103

0.97

102 101 0.01

0.1

1

10

Power (mW)

Figure 6. Excitation dependent micro-photoluminescence (µ-PL) of InN nano-pyramidal-prisms on patterned GaN nanowires using three-step growth method at low temperature (~28 K). (a) Normalized µ-PL, (b) PL peak emission energy (Epeak), (c) full width at half maximum (FWHM), and (d) PL peak intensity as a function of excitation optical power. The inset of (b) shows the evolution of energy at half of peak intensity in the low (ELOW) and high (EHIGH) energy side.

However, relative to EHIGH, a very small change in ELOW is observed, indicating very small or negligible localized states and band-tailing effect. Fig. 6(d) shows the PL peak intensity as a function of excitation power. It scales approximately linearly with power. A more careful fitting


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reveals the dependence as ~ P0.97 at low excitation powers, which is consistent with free-to-

(a)

P = 0.53 mW 285 K 261 K 230 K 200 K 169 K 140 K 121 K 100 K 81 K 59 K 40 K 28 K

0.6

0.7

0.8

0.9

0.70

Peak Energy (eV)


bound pair recombination at low temperatures [26].

(b)

0.69

1.5 mW 5.3 mW Eg (0) = 0.692 eV

0.68 γ = 0.42 meV β = 580 K

0.66 0.65 0.64

γ = 0.37 meV β = 512 K

Eg (0) = 0.689 eV

0.67

Fit equation: 2 = 0 − +

0

50 100 150 200 250 300

Energy (eV)

E0 (meV)

~45.9 meV

10

~3.14 meV ~88 K


Temperature (K)

P = 5.3 mW

ln[I0 / I(T) - 1]

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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~ exp −

285 K

28 K

0.60

0.65

0.70

0.75

Energy (eV)

~kT

(c) 0

(d) 100

200

300 -1

1/kT (eV )

400

100

500

Temperature (K)

Figure 7. Temperature dependent micro-photoluminescence (µ-PL) of InN nano-pyramidalprisms on patterned GaN nanowires using three-step growth method. (a) Normalized µ-PL, and (b) PL peak emission energy (Epeak) as function temperature. (c) ln[I0/I(T) - 1] as a function of 1/kT revealing two distinct PL quenching activation energy. (d) Specific energy, E0 as a function of temperature showing nearly kT dependence of E0. Inset show normalized PL for temperature ranging from 28 – 285 K.

Figure 7(a) shows normalized PL spectra varying temperature from 28K to 285 K at 5.3 mW of excitation power; spectra are vertically shifted for visual clarity. We observe a strong red shift of Epeak, accompanied by an increase in the FWHM with increasing temperature. Fig. 7(b) shows


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Epeak as a function of temperature for two different excitation powers (1.5 and 5.3 mW). The temperature dependence of Epeak is directly related to the temperature dependence of the bandgap energy and can be modeled using Varhni’s equation:

= 0 −

+

(1)

where and 0are bandgap at temperatures T and 0 K, respectively. Fitting Epeak from Fig. 7(b) with equation (1) gives 0 =0.689 and 0.692 eV, =0.42 and 0.37 meV, and =580 and 512 K for 1.5 and 5.3 mW of excitation power, respectively. These values are in close agreement with previous observations [22] [23] [27] [5] [28]. The temperature dependence of Epeak is not observed in n-type degenerate InN thin films or nanowires [24] [29] [30]. Observation of clear temperature dependence in Epeak further supports that the residual electron concentration in these nano-pyramidal-prisms is low. We also observe the quenching of PL peak intensity with temperature (I(T)) which can be modeled using the below equation [25],

=

1 + −

+ −

(2)

where, I0 is PL peak intensity at low temperature, and E1 and E2 are the activation energies related to different PL decay mechanisms. Shown in Fig. 7(c), we plot ln[I0/I(T) -1] as a function of 1/kT at an excitation power of 5.3 mW. Fitting using equation (2) reveals two distinct activation energies of ~45.9 meV and ~3.14 meV, which dominate at high (1/kT < 100 eV-1) and low (1/kT > 150 eV-1) temperatures, respectively. It is also clear that the change in decay mechanism occurs at ~88 K, which is in close agreement with previous observation of such a change at ~100 K [23]. This switching temperature (88K→7.6 meV) is known to be related to excitation binding energy [31] [32] [33] and is in close agreement with previously observed value of ~8 meV [23]. Furthermore, the excitation binding energy can be extracted from the


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activation energy related to PL decay at low temperature. We found this to be 3.14 meV for the nano-pyramidal-prisms in this study. Our observation is very close to previously reported experimental values and theoretical predictions in the range of ~2-3 meV [23] [34] [35] [36] [37]The normalized PL spectra (without the vertical shift) as a function of temperature are shown in the inset of Fig. 7(d). It is observed that the high energy emission decays slowly as the temperature increases. This decay can be modeled using ~exp(-E/E0), where E is the photon energy and E0 is specific energy related to the decay [4]. The solid bold line shows the fitting (with ~exp(-E/E0)) for the PL spectra taken at 285 K. Fig. 7(d) shows a nearly linear dependence of E0 as a function of temperature. The extension of the high energy tail represents the thermal distribution of photo-excited carriers in the energy band, which is expected to be related to the thermal energy kT. It can be seen that the E0 approximately follows the solid red line in Fig. 7(d) representing kT. , as expected. Such a dependence is only observed above 100 K in degenerate InN nanowires [4]. We observe such dependence down to low temperature (~28.2 K, the limit of our experimental setup) which indicates very high quality and uniform properties of these patterned InN nano-pyramidal-prisms. This also suggests very low or negligible Landsberg broadening [38] in contrast to degenerate InN nanowires [24], and confirms the presence of low residual carrier concentration in these nano-pyramidal-prisms. We also observe a secondary peak about 70 – 73 meV red shifted from the dominant PL peak which is attributed to the phonon sideband emission in these InN nano-pyramidal-prisms. The observation of such side band emission up to 261K is consistent with previous observations in InN nanowires with low residual electron concentration [23] [25] [28]. In summary, the growth of InN on patterned GaN nanowire templates was explored. Through the employment of a three-step growth technique, uniform hexagonal nano-pyramidal-prisms of


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InN were formed. Temperature dependent µ-PL showed the high quality, and intrinsic nature of the InN prisms. This opens the door to the formation of coalesced thin films of InN.

Table 1. Summary of the samples Sample name

Total Active Growth Time (hrs)

Sample A

3 hours growth at 425 °C

Sample B

3 hours growth at 450 °C

Sample C

Sample A followed by an 460 °C anneal for 20 min in N-plasma

Sample D

Sample C followed by a 0.5 hours InN growth at 500 °C

Corresponding Author *E-mail: [email protected].

ACKNOWLEDGMENT This work was supported by the National Science Foundation CAREER Award (DMR1055164). G.T.W. acknowledges funding from Sandia’s LDRD Program. B.L. acknowledges funding from the U. S. DOE, Office of Science, Office of Basic Energy Sciences, Materials


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Science and Engineering Division. Sandia National Laboratories is a multimission laboratory managed and operated by National Technology and Engineering Solutions of Sandia, LLC., a wholly owned subsidiary of Honeywell International, Inc., for the U.S. Department of Energy’s National Nuclear Security Administration under contract DE-NA0003525. REFERENCES [1] H. Lu, W. Schaff, and L. Eastman, "Surface chemical modification of InN for sensor applications," Journal of Applied Physics, 96, 3577 (2004). [2] Davydov, V.Yu., Klochikhin, A.A., Emtsev, V.V., Kurdyukov, D.A., Ivanov, S.V., Vekshin, V.A., Bechstedt, F., Furthmüller, J., Aderhold, J., Graul, J., Mudryi, A.V., Harima, H., Hashimoto, A., Yamamoto, A. and Haller, "Band Gap of Hexagonal InN," Phys. stat. sol. (b), 234, 787 (2002). [3] J. Wu, W. Walukiewicz, K. M. Yu, W. Shan, and J. W. Ager III, "Superior radiation resistance of In1-xGaxNIn1-xGaxN alloys: Full-solar-spectrum photovoltaic material system," Journal of Applied Physics, 94, 6477 (2003). [4] J. Wu, W. Walukiewicz, K.M. Yu, J.W. Ager III, E.E. Haller, H. Lu, W.J. Schaff, "Small band gap bowing in In1−xGaxN alloys," Appl. Phys. Lett., 80, 4741 (2002). [5] J. Wu, W. Walukiewicz, W. Shan, K. M. Yu, J. W. A. Iii, S. X. Li, E. E. Haller, H. Lu, and W. J. Schaff, "Temperature dependence of the fundamental band gap of InN," J. Appl. Phys., 94, 4457 (2003). [6] D. C. Look, H. Lu, W. J. Schaff, J. Jasinski, and Z. Liliental-Weber, "Donor and acceptor concentrations in degenerate InN," Appl. Phys. Lett., 80, 258 (2002). [7] V. W. L. Chin, T. L. Tansley, and T. Osotchan, "Electron mobilities in gallium, indium, and aluminum nitrides," Journal of Applied Physics, 75, 7365 (1994). [8] J. Wu, W. Walukiewicz, S. X. Li, R. Armitage, J. C. Ho, E. R. Weber, E. E. Haller, Hai Lu, William J. Schaff, A. Barcz, and R. Jakiela, "Effects of electron concentration on the optical absorption edge of InN," Appl. Phys. Lett., 84, 2805 (2004). [9] T. L. Tansley, and C. P. Foley, "Electron mobility in indium nitride," Electronics Letters, 20, 1066 (1984). [10] A. Kraus, C. Hein, H. Bremers, U. Rossow, and A. Hangleiter, "Growth kinetics and island evolution during double-pulsed molecular beam epitaxy of InN," Journal of Applied Physics, 119, 235308 (2016). [11] Y. F. Ng, Y. G. Cao, M. H. Xie, X. L. Wang, and S. Y. Tong, "Growth mode and strain evolution during InN growth on GaN(0001) by molecular-beam epitaxy," Appl. Phys. Lett., 81, 3960 (2002). [12] T. Stoica, R. Meijers, R. Calarco, T. Richter, and H. Luth, "MBE growth optimization of InN nanowires," J. Cryst. Growth, 290 (2006), p. 241. [13] A. T. M. Golam Sarwar, S. D. Carnevale, T. F. Kent, M. R. Laskar, B. J. May, and R. C. Myers, "Molecular beam epitaxy of InN nanowires on Si," J. Cryst. Growth, 428, 59


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For Table of Contents Use Only Title: "Hexagonal nano-pyramidal-prisms of nearly intrinsic InN on patterned GaN nanowire arrays"

GaN

PL

Author(s): Golam Sarwar, A.T.M.; Leung, Benjamin; Wang, George; Myers, Roberto

T

0.60

InN

Eg (eV)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


0.65 0.70 E (eV)

0.68 0.66 0

100 200

T (K)

SYNOPSIS By separating nucleation and growth processes, hexagonal nano-prisms of InN are grown by molecular beam epitaxy on top-down patterned GaN nanowires. Subsequent growth of InN on already formed InN nano-prisms at elevated temperatures leads to high optical quality, nearly intrinsic single crystalline InN nano-prism arrays with pyramidal tops. Coalescence of these ordered arrays could be used to obtain free-standing InN.


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Hexagonal Nanopyramidal Prisms of Nearly Intrinsic InN on Patterned

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