InGaN multi-quantum-well coaxial nanorods

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ScienceDirect Acta Materialia 65 (2014) 118–124 www.elsevier.com/locate/actamat

Ultrashort carrier lifetime of vapor–liquid–solid-grown GaN/InGaN multi-quantum-well coaxial nanorods Mohamed Ebaid a, Jin-Ho Kang a, Seung-Hyuk Lim b, Suk-Min Ko b, Yong-Hoon Cho b, Sang-Wan Ryu a,⇑ b

a Department of Physics, Chonnam National University, Gwangju 500-757, Republic of Korea Department of Physics and KI for the NanoCentury, Korea Advanced Institute of Science and Technology, Daejeon 305-701, Republic of Korea

Received 14 October 2013; received in revised form 18 November 2013; accepted 19 November 2013 Available online 27 December 2013

Abstract Luminescence and carrier dynamics of GaN/InGaN multi-quantum-well coaxial nanorods (MCNRs) were studied by means of photoluminescence (PL), cathodoluminescence (CL) and time-resolved PL (TRPL). The PL of as-grown MCNRs showed an intense blue emission together with broad emission at longer wavelengths. CL measurements of several single MCNRs attributed the broad emission to the wetting layer frequently observed in vapor–liquid–solid-grown nanorods. Non-single exponential intensity decays were observed by TRPL, which were ascribed to the In fluctuation in the InGaN alloy. Radiative and non-radiative lifetimes were then calculated via a stretched exponential model. An ultrafast carrier lifetime in the range of a few tens of picoseconds along with a high internal quantum efficiency (IQE) of about 59% resulted. The ultrafast carrier lifetimes were attributed to the improvement in the carrier collection efficiency due to the radial heterostructuring of GaN with InGaN shells, while the high IQE implied that carriers were mostly recombined radiatively. This study reveals that the coaxial growth of InGaN with GaN nanorods resulted in an ultrafast carrier lifetime and a luminescence efficiency that could be controlled by adjusting the growth temperature gap between the GaN and the InGaN. Ó 2013 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: GaN; Nanowires; Time-resolved photoluminescence; Lifetime

1. Introduction GaN/InGaN multi-quantum-well coaxial nanorods (MCNRs) have been developed as versatile nanoscale materials for visible light emission to replace the conventional planar multi-quantum wells (MQWs) in optoelectronic device applications [1–7]. In the context of device efficiency, various inherent problems related to the growth of MQWs on bulk GaN have been disabled by employing dislocation-free GaN nanorods as growth templates [3,8]. For example, nanometer-scale GaN nanorods may hinder the formation of extended defects such as dislocations, stacking faults and twins [3,9]. Consequently, smooth growth interfaces between highly mismatched materials ⇑ Corresponding author. Tel.: +82 62 530 3476; fax: +82 62 530 3369.

E-mail address: [email protected] (S.-W. Ryu).

such as GaN and InGaN can be established, which may significantly reduce optical losses at the nanorod–air interface [7]. Furthermore, the three-dimensional geometry of nanorods enables the growth of MQWs on planes other than the c-plane, which in turn may reduce the large strain-induced built-in piezoelectric and polarization fields [2,3,5,7]. InGaN MQWs embedded in the axial direction [10,11] or grown coaxially with GaN nanorods [3,5,7] have previously been demonstrated. GaN/InGaN MCNRs were fabricated by various techniques, such as the post-growth of MQWs on the surface of etched GaN pillars [2], selective area growth [3] and the vapor–liquid–solid (VLS) growth technique [7,12]. VLS may be considered the simplest way to grow GaN nanorods; however, radial growth of GaN/InGaN by this mechanism may be more challenging due to the complex growth conditions required to

1359-6454/$36.00 Ó 2013 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.actamat.2013.11.058

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overcome the great competition between side wall nucleation and direct reaction with metal catalysts [7]. For instance, to fully cover GaN cores with continuous InGaN shells, two prerequisites must be satisfied: a low growth temperature [2,3,5] and nitrogen-rich conditions [7,12]. The former is mainly adopted to reduce In desorption and the latter to compensate the incomplete pyrolysis of NH3 at low temperatures. In contrast, high quality GaN requires high growth temperatures and hydrogen carriers. Therefore, the growth of InGaN on GaN template may be limited by these compromises. In particular, the interfaces between GaN and InGaN may be degraded due to the associated large temperature gap between their successive growths [13–15]. Therefore, the growth temperature gap between GaN and InGaN can be considered a critical parameter that may strongly influence the optical characterization of these heterostructures. In particular, the growth temperature of InGaN can significantly affect the carrier transportations due to the correlated influence on the In concentration and the interface flatness between barrier and well layers [15]. Furthermore, the carrier dynamics may be influenced by the structural geometry of the desired nanorod heterostructures. For instance, carrier recombination may be accelerated in a core/shell geometry compared to that in axially grown nanorods due to the efficient radial confinement and the improved carrier collection efficiencies in the former design [16,17]. To the best of our knowledge, the carrier dynamics of GaN/InGaN MCNRs have not yet been studied, though those of the axial GaN/InGaN nanorods have been reported in a number of papers [11,18–21]. Recently, we fabricated high-quality GaN nanorods with a high aspect ratio by the VLS growth mechanism at a relatively low growth temperature of 765 °C by employing a new Ni/In/Ga ternary alloy catalyst with a low melting point [22]. This temperature is much lower than that used in metal–organic chemical vapor deposition (MOCVD) and may be comparable to that of molecular beam epitaxy (MBE). At such a low growth temperature of GaN cores, high-quality InGaN shells may be formed due to the corresponding small temperature gap. Decreasing the temperature gap will also decrease the effective annealing time of the InGaN alloy, leading to MQW structures with higher In fractions [15]. The effect of the InGaN growth temperature on the luminescence and carrier dynamics of GaN/InGaN MCNRs are addressed in detail in this study. 2. Experimental method GaN/InGaN MCNRs were grown in a vertical MOCVD reactor by the VLS growth mechanism. Low-temperature growth of GaN cores was accomplished by designing a new catalyst based on the insertion of an In interlayer between a very thin Ni-metal film and a Ga-capping layer. GaN cores were synthesized on c-plane sapphire substrates at 765 °C and 75 torr for 30 min. Trimethylgallium (TMGa) and NH3 were used as Ga- and N-precursors,

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respectively, while H2 was used as the carrier gas. The flow rates of TMGa and NH3 were fixed at 8.8 lmol min1 and 13.4 mmol min1, respectively. Under these conditions, quasi-aligned GaN cores were formed having lengths up to 8 lm and diameters of about 150 ± 30 nm. After the formation of these high aspect ratio GaN cores, the growth conditions were adjusted to support deposition on side walls and prevent the direct nucleation of reactants with the catalyst. During the growth of shell layers, TMGa, trimethylindium (TMIn) and NH3 were introduced continuously into the MOCVD reactor in the presence of N-carrier. The flow rates of TMGa and NH3 were fixed at 4.7 lmol min1 and 67 mmol min1, respectively, while that of TMIn was changed between 1.3, 2.3 and 3.2 lmol min1. The reactor pressure was maintained at 75 torr. While the growth temperature of the InGaN quantum well was varied between 715 and 745 °C, GaN barriers were grown at the same temperature as the GaN cores at 765 °C. NH3 flowed continuously during the temperature change between successive growths of GaN and InGaN to avoid surface degradation. The morphological properties of GaN/InGaN MCNRs were studied by field emission scanning electron microscopy (FE-SEM; JSM-6700F, JEOL Japan). The luminescence of the as-grown MCNRs was characterized via room-temperature photoluminescence (PL) using an He– Cd laser (325 nm) with an optical power of about 20 mW. Local emission of a single GaN/InGaN MCNR was measured from spatially resolved cathodoluminescence (CL) (Hitachi 4300SE) with a step size of about 2 nm. Time-resolved PL (TRPL) was used to study carrier dynamics as a function of InGaN growth temperature. During TRPL measurements, the as-grown samples were excited by a Ti:sapphire pulsed laser at a fixed optical power of about 10 lW. The time resolution was less than 300 fs and the repetition rate was about 4 MHz. The excitation wavelength of the pulsed laser was adjusted to 400 nm to excite the InGaN quantum well (QW) only. All TRPL measurements were conducted at room temperature. 3. Results and discussion 3.1. Morphology of GaN/InGaN MCNRs Fig. 1 shows typical FE-SEM images of the morphology of GaN/InGaN MCNRs. The magnified image of a single GaN/InGaN MCNR shown in Fig. 1b reveals that these nanorods have an isosceles triangular cross-section with smooth side walls. This kind of triangular cross-section was previously reported for GaN/InGaN MCNRs grown by MOCVD via the VLS technique under similar conditions [7,23,24]. In these reports, the two congruent sides of the triangular cross-section were indexed to f1  1 0 1g semi-polar facets with a c-plane base. Similarly, high-quality GaN bulk grown by epitaxial lateral overgrowth also had a triangular cross-section with f1 1 0 1g facets, which

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number of QWs [27], only two periods of MQW were designed. PL spectra are presented in Fig. 2a, where a temperature-sensitive luminescence was observed within a very small temperature window of 30 °C. With increasing growth temperature of InGaN from 715 to 730 °C, a blue shift of the main emission peak by about 12 nm was observed, indicating a significant decrease in In concentration in the InGaN alloy. The gradual increase in temperature to 745 °C also led to a blue shift of the main peak by about 4.5 nm. Further, PL spectra revealed the presence of a broad emission at longer wavelengths, which almost vanished at 730 °C. This kind of broad emission was previously observed for GaN nanorods covered with MQWs grown on non-patterned substrates [5,14]. The origin of this emission was attributed to either the native structural and chemical defects [5] or the inhomogeneous incorporation of In due to the random spontaneous nucleation on such substrates [14,28]. Interestingly, this broad emission was also dependent on the density of GaN/InGaN MCNRs, as shown in Fig. 2b. To explain this observation, we should consider the effects related to the interfacial parasitic layer or the wetting layer, which inherently evolved during the growth of nanorods on non-patterned substrates [29,30]. The relative emission of the blue to yellow peaks reached a maximum for the highest nanorod density obtained at 730 °C. This may be attributed to the full Fig. 1. (a) FE-SEM image of as-grown GaN/InGaN MCNRs and (b) magnified image of a single GaN/InGaN MCNR; inset is a cross-sectional view.

360

400

440

480

520

560

600

Wavelength (nm)

Iblue/ Iyellow

30

I

(b)

blue

0.88

/I

yellow

Nanorod density

25

0.80

20

00.27

15

0.64

10

0.56 0.48

5 0

Nanorods Density (cm-3)

The optical emission of the as-grown GaN/InGaN MCNRs was characterized by using room-temperature PL. We emphasize that, during the growth of these nanorods, the GaN quantum barrier (QB) was grown at the same temperature as the GaN core (765 °C), with only the temperature of the InGaN QW being changed between 715 to 745 °C. In order to clearly correlate between the InGaN growth temperature and the luminescence of the GaN/InGaN MCNRs, and to avoid the compositional modulation reported for the InGaN alloy with increasing

715 C o 730 C o 745 C

PL Intensity (a.u.)

was attributed to the thermodynamic stability of these planes under a wide range of growth conditions [25]. In our case, the two congruent sides of the triangular crosssection made an angle of about 58° with the base, which was measured from a cross-sectional view SEM image, as shown in the inset of Fig. 1b. By comparing our observations with previous reports of similar nanorods, as well as with bulk GaN and based on the calculated angles, we may suggest that our MCNRs were also grown with f1 1 0 1g facets and a (0 0 0 1) base. Consequently, the primary axis of the GaN/InGaN MCNRs is along the h1 1  2 0i direction [23,26]. 3.2. Room-temperature photoluminescence

o

(a)

0.40 715

720

725

730

735

740

745

InGaN Growth Temperature (oC) Fig. 2. (a) Room-temperature PL of the as-grown GaN/InGaN MCNRs as a function of InGaN growth temperature and (b) shows the variation of the relative emission between blue and broad band peaks and the corresponding nanorods density.

M. Ebaid et al. / Acta Materialia 65 (2014) 118–124

coverage of the substrate, which can significantly screen the luminescence from the wetting layer. This behavior suggested that the observed broad emission could be originated from the wetting layer and was not the luminescence from GaN/InGaN MCNRs. 3.3. Characterization of single GaN/InGaN MCNR by cathodoluminescence In order to clarify the results suggested by room-temperature PL, we studied the local emission of several single MCNRs with their InGaN QW grown at 730 °C by means of CL, as presented in Fig. 3. To be measured by CL measurements, as-grown GaN/InGaN MCNRs were first detached from the sapphire substrate and then dispersed on a clean Si substrate. A sectional SEM image of a scanned nanorod was first taken, and then the detection mode was switched to measure the integrated emission spectra and to record the spectrally resolved CL images under the same magnification. Contrary to the PL measurements collected from GaN/InGaN MCNRs ensembles, only an intense blue emission of the MQW was observed for all positions of the single nanorod. Furthermore, the spectrally resolved CL images showed that the yellow emission from the single nanorod was faint and negligible compared to that collected at the blue wavelength, as shown in Fig. 3b. Spatially resolved CL results may be considered clear evidence confirming that the yellow emission was

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originated from the wetting layer formed unintentionally during the growth of GaN cores and not from the nanorod. In addition, no significant emission for the band-to-band transition of GaN cores was observed, suggesting a full radial coverage of the GaN cores with GaN/InGaN MQW shells. It is worth mentioning that a strong and broad defect-related emission was previously observed in the CL spectra of single GaN/InGaN core/shell nanorods with similar triangular cross-sections [31]. These nanorods were grown by the Ni-assisted VLS mechanism, but with a growth temperature gap of about 190 °C between the GaN core and the InGaN shell, compared to one of only 35 °C in our case. 3.4. Tunable emission of GaN/InGaN MCNRs Since the growth of InGaN at 730 °C showed no significant defect-related yellow luminescence, the controllability of the optical emission by changing the molar flow ratio between TMIn and TMGa was investigated at the same temperature. During this experiment, the NH3 flow rate was fixed at 67 mmol min1, while the molar flow ratio (TMIn/TMIn + TMGa) was changed between 0.22, 0.33 and 0.41. As shown in Fig. 4a, the PL emission red shifted with increasing molar flow ratio – clear evidence of enhanced In incorporation. The red shift attained its maximum of about 467 nm at the highest In to Ga ratio of 0.41

(a) Normalized CL Intensity (arb. unit.)

Apex(1) Stem (2) Root (3)

Normalized PL Intensity (arb.unit.)

(a)

350 300

350

400

450

500

550

600

0.22 0.33 0.41

400

450

500

550

600

Wavelength (nm)

650

Wavelength (nm) 480

Wavelength FWHM

(b)

55 460 450

50

440

FWHM (nm)

Wavelength (nm)

470

60

45 430 420 0.20

0.25

0.30

0.35

0.40

40

TMIn/TMIn+TMGa Fig. 3. (a) CL spectra from several detection points of a single GaN/ InGaN MCNRs and (b) the associated spectrally resolved CL images of the same nanorod, which were collected at blue and yellow emissions.

Fig. 4. (a) Controllable emission of GaN/InGaN MCNRs by changing the molar flow ratio of In/Ga precursors and (b) the measured wavelength shift and FWHM as the ratio of the In/Ga precursors.

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without the appearance of defect-related yellow luminescence. The nominal change of the full width at half maximum (FWHM) calculated with increasing In concentration is shown in Fig. 4b, and indicates good optical and structural qualities. In contrast, band broadening combined with defect-related emission has been frequently reported for either MQW nanorods or planar MQWs with increasing In concentration [3,5,15]. Essentially, achieving tunable emission with good structural and optical qualities may suggest these MCNRs as potential candidates in device applications. 3.5. Carrier dynamics of GaN/InGaN MCNRs The luminescence of GaN/InGaN MCNRs was further studied by TRPL to trace carrier recombination dynamics in the InGaN QW according to the change of its growth temperature. 3.5.1. Temporal decay maps An overall view of the carrier dynamics can be given by the temporal decay maps measured by recording the PL intensity decay against the wavelength. These are depicted in Fig. 5. It should be noted that the PL decay is very sensitive to change in the InGaN QW growth temperature, i.e. to the temperature gap between QB and QW. At the lowest growth temperature, the PL intensity showed a broad decay spectrum and shifted to longer wavelengths. This broad spectrum, as well as the PL red shift, was previously attributed to the recombination of localized excitons at the potential minima or to the clusters of higher In fraction formed due to In composition fluctuations within the InGaN alloy [32–34]. Since the samples were grown at different temperatures and the In desorption rate was strongly influenced by the growth temperature, it was then expected that the In concentration would deviate between the samples. As a result, InGaN grown at 715 °C is expected to contain more In than those grown at higher growth temperatures, which may increase the chance for more In fluctuations. Therefore, as the QW was excited, carriers were diffused towards the lowest energy states and the process is accompanied by a significant red shift of the PL [34,35]. As the growth temperature was increased to 730 °C, the effects related to In fluctuations may be less

pronounced and a faster PL decay, with a blue shift and a narrower spectrum, was observed. A gradual increase in temperature to 745 °C led to another slight red shift, which may be attributed to the randomness of the desorption rate of In [15]. This randomness of the In desorption process may cause non-uniform In distribution in the InGaN QW, i.e. the existence of some localized states, which can be responsible for the observed small red shift.

3.5.2. Stretched exponential modeling and carrier lifetimes calculation It was reported that In concentration fluctuations in the InGaN QW may cause a stretched exponential decay [32,36]. In such decay profiles the determination of recombination lifetimes is not straightforward due to the additional complication arising from the non-single exponential decay. Consequently, the TRPL decay spectra were fitted by the stretched exponential model, which can be expressed as: b

IðtÞ ¼ I 0 exp½ðt=tÞ 

ð1Þ

where I(t) is the PL intensity at time t, I0 is the initial PL intensity, b is the stretching parameter related to the localized states in the QW and s is the time constant [21,32,36,37]. This model was widely used to describe the PL decay of disordered structures. Fig. 6 shows the temporal decay profiles as a function of InGaN growth temperature and independent calculations of b and s values by plotting the double logarithm of the PL decay vs. the natural logarithm of time. As shown in Table 1, the b-value was almost unity at the growth temperature of 730 °C, which may indicate a small In fluctuation, while its value was reduced at the other growth temperatures. Thus, the PL decay at 730 °C may be defined as a single exponential decay [37]. On the other hand, the values calculated for the time constant s revealed ultrafast carrier recombination lifetimes within a few tens of picoseconds, which was slightly reduced with increasing growth temperature of InGaN. The radiative (sr) and non-radiative (snr) lifetime components can be derived by combining the equations of internal quantum efficiency (IQE) and initial lifetime (s) [21,37], which may be given respectively by:

Decay Time (ns)

350.0

0.1

0.1

0.1

136.0

0.2

0.2

0.2

78.00

0.3

0.3

0.3

36.00

0.4

0.4

(a)

0.5 400

715 oC 450

500

Wavelength (nm)

550

0.4

(b)

0.5 400

730 oC 450

500

Wavelength (nm)

550

10.00

(c)

0.5 400

o

745 C 450

500

0

550

Wavelength (nm)

Fig. 5. Temporal decay maps of GaN/InGaN MCNRs as a function of InGaN growth temperature measured by TRPL, (a) 715 °C, (b) 730 °C and (c) 745 °C.

Ln [Ln(I(o)/I(t))] Ln [Ln(I(o)/I(t))] Ln [Ln(I(o)/I(t))] (arb. unit.) (arb. unit.) (arb. unit.)

M. Ebaid et al. / Acta Materialia 65 (2014) 118–124

715

(a)

730

PL Intensity (arb. unit.)

745

0.0

0.2

0.4

0.6

0.8

1.0

2.0 1.5

(b)

1.0 0.5

715 oC

0.0 1.8

(c)

1.2 0.6 0.0 1.5

730 oC

(d)

1.0 0.5 0.0

745 oC

-3.5 -3.0 -2.5 -2.0 -1.5 -1.0

Ln [Time(ns)]

Decay Time (ns)

Fig. 6. (a) Temporal decay curves of GaN/InGaN MCNRs as a function of InGaN growth temperature measured by TRPL and (b)–(d) are the corresponding fittings by stretched exponential function.

IQE ¼ ð1 þ sr =snr Þ

1

ð2Þ

1=s ¼ 1=sr þ 1=snr

ð3Þ

The s values were obtained from the stretched exponential fit, while that of the IQE were measured from the PL spectra shown in Fig. 7 as the ratio between the PL peak intensity at room temperature to that measured at 15 K, assuming the freeze-out of defect state transitions at 15 K. As shown in Fig. 7, no defect-related luminescence appeared at 730 °C, which was reflected in the high IQE of 59% compared to those of 19 and 31% at 715 and 745 °C, respectively. The values of sr and snr were then calculated and listed in Table 1. Apparently, sr was strongly influenced by the growth temperature of InGaN, while snr changed only slightly with the temperature. For the sample grown at 730 °C, which showed the highest IQE

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and the least In fluctuation, sr was smaller than snr – clear evidence of its good quality. On the other hand, carriers needed a longer time to recombine radiatively at other growth temperatures. Indeed, by referring to the CL data given in Fig. 3, which indicates that there is no defect-related luminescence at 730 °C, and based on the calculated values for IQE and b of about 59% and 1, respectively, at this temperature, we may conclude that the observed ultrafast recombination lifetime of 24 ps is due to the efficient radiative recombination of photo-excited carriers in the InGaN QW. It is worth mentioning that the improved surface-tovolume ratio may cause the surface effects to play an important role in the recombination dynamics of nanorods. Surface defect states may remarkably increase the recombination lifetime and release the non-radiative surface recombination that has detrimental effects on the radiative efficiency [18,21]. For instance, recombination lifetimes of GaN/InGaN NRs fabricated by wet etching were in the range of several to tens of nanoseconds [18,20], compared to a lifetime of about 24 ps in our VLS-grown NRs. The long lives of the carriers observed in nanorods fabricated by etching are attributed to the surface damage during sample processing, which may result in the spatial separation of carriers due to band bending near the surface of nanorods. Consequently, much effort has been devoted to the passivation of the surface of GaN NRs in order to enhance their emission efficiency [21,38,39]. GaN/InGaN axial nanorods grown by Ga-assisted MBE showed carrier recombination rates in the range of several hundreds of picoseconds after the passivation of their surfaces with silicon nitride or parylene [21]. The IQE of these nanorods was also measured and reached a maximum at 52%, compared to 59% for our optimized InGaN growth temperature. The faster recombination rates of MCNRs than of axially grown nanorods or those made by

Table 1 Carrier lifetimes as calculated from TRPL decay in terms of the growth temperature of InGaN. Data were analyzed by using the stretched exponential model. Stretching parameter (b)

Time constant, s (ps)

Internal quantum efficiency, IQE (%)

Radiative lifetime, sr (ps)

Non-radiative lifetime, snr (ps)

715 730 745

0.72 0.99 0.82

35 24 21

19 59 31

184 40.7 67.7

43.2 58.5 30.4

PL Intensity (arb. unit.)

InGaN growth temperature (°C)

6000 5000

10000

(a)

15 K 300 K 715 oC

4000

8000

(b)

4000

730 oC

6000

3000

5000

15 K 300 K

(c)

15 K 300 K 745 oC

3000

4000

2000

2000

1000

2000 1000 0 350

400

450

500

Wavelength (nm)

550

0 350

400

450

500

Wavelength (nm)

550

0 350

400

450

500

550

Wavelength (nm)

Fig. 7. PL spectra measured at room temperature (300 K) and 15 K of GaN/InGaN MCNRs as a function of InGaN growth temperature, (a) 715 °C, (b) 730 °C and (c) 745 °C. IQE was calculated as the ratio between PL peak intensity at 300 K to that of 15 K.

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wet etching could be attributed to the enhanced carrier collection efficiency of the coaxial QW structure due to the small distance the carriers would travel before capture and the absence of the piezoelectric field associated with the growth in non-polar growth directions [39]. Moreover, as compared to the post passivation of NRs that may cause additional surface damages, the density of surface states may be effectively reduced by the in situ epitaxial deposition of GaN/InGaN MQW shells after the growth of GaN NRs. From the values listed in Table 1, we may conclude that the growth temperature of InGaN QW and hence the temperature gap between the QB and QW has a substantial effect on the carrier dynamics in MCNRs, which can be optimized to achieve ultrafast recombination without affecting the IQE. In our case, we found that a growth temperature gap of only 35 °C can induce ultrafast recombination without degradation of the radiative efficiency, which may be attributed to the associated improvement of crystal quality [15]. 4. Summary and conclusion In conclusion, the luminescence and carrier dynamics of GaN/InGaN MCNRs were studied by PL, CL and TRPL. There was a strong relationship between luminescence and carrier dynamics, and the growth temperature gap between QB and QW. CL measurements of several single nanorods displayed very slight defect-related yellow luminescence from GaN/InGaN MCNRs. Non-single exponential decay of PL intensity was recorded by TRPL, which was attributed to In fluctuations in the InGaN alloy. Because of the efficient capture of photogenerated carriers and the small piezoelectric field in the h1 1  2 0i growth direction, an ultrafast recombination lifetime was obtained together with a high IQE. These findings are extremely important, as they can pave the way for the fabrication of high-quality light nano-emitters. Acknowledgements This work was supported by the Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education, Science and Technology (2011-0017190). References [1] Shen G, Chen D. Optoelectron Chin 2010;3:125. [2] Chang JR, Chang SP, Li YJ, Cheng YJ, Sou KP. Appl Phys Lett 2012;100:261103. [3] Yeh TW, Lin YT, Stewart LS, Dapkus PD, Sarkissian R, Brien JDO, et al. Nano Lett 2012;12:3257. [4] Guo W, Banerjee A, Bhattacharya P, Ooi BS. Appl Phys Lett 2011;98:193102. [5] Koester R, Hwang JS, Salomon D, Chen X, Bougerol C, Barnes JP, et al. Nano Lett 2011;11:4839.

[6] Lin HW, Lu YJ, Chen HY, Lee HM, Gwo S. Appl Phys Lett 2010;97:073101. [7] Qlan F, Li Y, Gradecak S, Park HG, Dong Y, Ding Y, et al. Nat Mater 2008;7:701. [8] Wu YR, Chiu C, Chang CY, Yu P, Kuo HC. IEEE J Sel Top Quant Electron 2009;15:1226. [9] Kuykendall T, Pauzauskie P, Lee S, Zhang Y, Goldberger J, Yang P. Nano Lett 2003;3:1063. [10] Tourbot G, Bougerol C, Glas F, Zagonel LF, Mahfoud Z, Meuret S, et al. Nanotechnology 2012;23:135703. [11] Banerjee A, Dogan F, Heo J, Manchon A, Guo W, Bhattacharya P. Nano Lett 2011;11:5396. [12] Kuykendall T, Aloni S, Plante IJL, Mokari T. Int J Photoenergy 2009;2009:1. [13] Albert S, Encabo AB, Lefebvre P, Barbagini F, Garcia MAS, Calleja E, et al. Appl Phys Lett 2012;100:231906. [14] Limbach F, Gotschke T, Stoica T, Calarco R, Sutter E, Ciston J, et al. J Appl Phys 2011;109:014309. [15] Olaizola SM, Pendlebury ST, Neill JPO, Mowbray DJ, Cullis AG, Skolnick MS, et al. J Phys D Appl Phys 2002;35:1. [16] Li Y, Qian F, Xiang J, Lieber CM. Mater Today 2006;10:18. [17] Hiruma K, Tomioka K, Mohan P, Yang L, Noborisaka J, Hua B, et al. J Nanotechnol 2012;2012:1. [18] Chan CCS, Reid BPL, Taylor RA, Zhuang Y, Shields PA, Allsopp DWE, et al. Appl Phys Lett 2013;102:111906. [19] Cardin V, Bertrand LID, Gre´goire P, Nguyen HPT, Sakowicz M, Mi Z, et al. Nanotechnology 2013;24:045702. [20] Jiang B, Zhang C, Wang X, Xue F, Park MJ, Kwak JS, et al. Opt Exp 2012;20:13478. [21] Jahangir S, Mandl M, Strassburg M, Bhattachary P. Appl Phys Lett 2013;102:071101. [22] Ebaid M, Kang J, Lee JK, Ryu SW. J Phys D Appl Phys 2013;46:385105. [23] Qian F, Li Y, Gradecak S, Wang D, Barrelet CJ, Lieber CM. Nano Lett 2004;4:1975. [24] Wang GT, Talin AA, Werder DJ, Creighton JR, Lai E, Anderson RJ, et al. Nanotechnology 2006;17:5773. [25] Hiramatsu K. J Phys Condens Matter 2001;13:6961. [26] Romanov AE, Baker TJ, Nakamura S, Speck JS. J Appl Phys 2006;100:023522. [27] Minsky MS, Fleischer SB, Abare AC, Bowers JE, Hu EL, Keller S, et al. Appl Phys Lett 1998;72:1066. [28] Tourbot G, Bougerol C, Grenier A, Hertog MD, Giao DS, Cooper D, et al. Nanotechnology 2011;22:075601. [29] Purushothaman V, Ramakrishnan V, Jeganathan K. RSC Adv 2012;2:4802. [30] Li Q, Wang GT. Appl Phys Lett 2008;93:043119. [31] Baird L, Ong CP, Cole RA, Haegel NM, Talin AA, Li Q, et al. Appl Phys Lett 2011;98:132104. [32] You G, Guo W, Zhang C, Bhattachary P, Henderson R, Xu J. Appl Phys Lett 2013;102:091105. [33] Hsu WT, Liao YA, Lu SK, Cheng SJ, Chiu PC, Chyi JI, et al. Physica E 2010;42:2524. [34] Krestnikov IL, Ledentsov NN, Hoffmann A, Bimberg D, Sakharov AV, Lundin WV, et al. Phys Rev B 2002;66:155310. [35] Krestnikov IL, Sakharov AV, Lundin WV, Usikov AS, Tsatsulnikov AF, Musikhin YG, et al. Phys Status Solidi (a) 2002;192:49. [36] Pophristic M, Long FH, Tran C, Ferguson IT, Karlicek RF. Appl Phys Lett 1998;73:3550. [37] Onuma T, Uchinuma Y, Suh EK, Lee HJ, Sota T, Chichibu SF. Jpn J Appl Phys 2003;42:L1369. [38] Armstrong A, Li Q, Lin Y, Talin AA, Wang GT. Appl Phys Lett 2010;96:163106. [39] Kar A, Li Q, Upadhya PC, Seo MA, Wright J, Luk TS, et al. Appl Phys Lett 2012;101:143104.