Partial edge dislocations comprised of metallic Ga bonds in

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Partial edge dislocations comprised of metallic Ga bonds in heteroepitaxial GaN Moonsang Lee, Hionsuck Baik, Wontaek Ryu, Yewon Jo, SeongYoung Kong, and mino yang Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.8b01488 • Publication Date (Web): 03 Jul 2018 Downloaded from http://pubs.acs.org on July 4, 2018

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Partial edge dislocations comprised of metallic Ga bonds in heteroepitaxial GaN Moonsang Lee†, Hionsuck Baik‡, Wontaek Ryu¶, Yewon Jo‡, SeongYoung Kong‡,

Mino Yang*‡,§

†

Korea Basic Science Institute, Daejeon 34133 Korea

‡Korea

¶Center

Basic Science Institute Seoul, Seoul 02841Korea

for Inter-University Research Facility, Kookmin University, Seoul 02707 Korea

§Laboratoire

de Physique des Solides (UMR CNRS 8502), Université Paris Sud, Orsay 91405, France

ABSTRACT

We investigated the atomic structure of inclined threading edge dislocation (TED) typically observed in GaN grown on Si(111) through (scanning) transmission electron microscopy. Atomic observations verified that the inclined TED consisted of two partial dislocations. These results imply that the inclined TED possesses a Ga-Ga atomic configuration that is

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energetically unfavorable. However, the introduction of such structures is considered unavoidable because the TEDs should climb regularly to mediate the applied stress or the increasing surface due to the buffer layer. This Ga-Ga configuration is highly likely to form metallic bonds and appears to be the primary reason for the inferior efficacy of a GaN lightemitting diode grown on Si(111).

KEYWORDS: partial edge dislocation, GaN, optical sectioning, metallic bonding

To improve the luminous efficacy of wurtzite GaN devices, dislocations should be eliminated or transformed into different species with fewer electronic defects.1 Current issues with regard to GaN (3.4 eV direct bandgap) involve the use of silicon substrates to replace sapphire for significant economic benefits.2 To obtain a flat, wide, and non-cracked GaN film on Si (GaN/Si), thermal mismatch between two materials should be controlled with buffer layers, such as AlxGa1-xN.3,4 The highly convex GaN/Si become flat at room temperature by the buffer layer. The coefficients of thermal expansion (10−6/K−1) are 7.5, 2.6, and 5.6 for sapphire, silicon, and GaN, respectively. (In this report, GaN/Si includes the buffer layers.) The excellent luminous efficacy of GaN even with a high density of dislocations is attributed to the atomic structures of threading edge dislocations (TED), which are electrically inert.5 Dislocations in GaN/Si have been observed to incline significantly in many published reports, unlike TEDs in GaN grown on sapphire (GaN/sapphire) lying parallel to the c-axis.6–9 The inclination is due to the unique growth condition, which is the highly curved wafer placed under severe compressive stress by the buffer layers.3 The efficacy of GaN/Si light-emitting diode is significantly lower than that of GaN/sapphire (up to ~70%). The inferiority has been known to be attributed to high dislocation density and reduced light extraction efficacy of GaN/Si.2 However, when TEDs incline, they experience a change in atomic structure. The


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inclination probably switches the TEDs to an electrically active state, thereby affecting efficacy. The examined GaN was grown on a Si(111) wafer using a commercial metal organic chemical vapor deposition (MOCVD) reactor (Veeco). The Si substrates were immersed in H2SO4:H2O2:H2O (3:1:1) and HF (6%) solutions to yield a hydrogen-terminated surface and a native oxide-free Si substrate, respectively. An ~150-nm-thick AlN layer and an ~250-nmthick Al0.4Ga0.6N buffer layer were grown sequentially on Si substrate at 1000 °C. An ~2-µmthick GaN layer was subsequently grown on top of these layers at 1040 °C.4 To characterize the GaN dislocations in plan- and cross-sectional views, surface relaxation contrast and weak beam dark field (WBDF) imaging techniques were employed.10,11 Both WBDF and surface relaxation contrast imaging methods utilize variations in the diffraction conditions around a dislocation core. Cross-sectional WBDF discriminates the types of dislocations when there is no contrast at g·b*u=0.12 Here, g, b, and u denote the diffraction spot, Burgers vector, and beam direction, respectively. The TED core exhibits no contrast at g=0001, because its b is normal to the core (b⊥core). By contrast, threading screw dislocations (TSDs) whose b value is parallel to the core (b||core) exhibit no contrast at g=1010. Surface relaxation contrast imaging was performed with plan-view bright field observation using g=2-1-10 and an 18° specimen tilt from [0001].11 In this technique, lines appear for TEDs, but present no contrast for TSD because b is parallel to the beam direction (b||u).12 Instead, TSDs are observed as spots because the atoms around the core emerging from the surface are rotationally relaxed, known as Eshelby twist.13 To observe the atomic structure changes from both views, we used optical sectioning techniques with a scanning transmission electron microscope (STEM). This technique acquires images with respect to specimen depth using a convergent beam. Varying the electron probe focus along the c-axis, atomic displacement relative to the a-axis of


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dislocation can be observed in plan-view specimens. The applied STEM (Titan, FEI) was equipped with a Schottky-type field emission gun, a high-angle annular dark field (HAADF) detector, and a probe Cs corrector (CESCOR, CEOS) enabling a spatial resolution below 1 Å. The calculated depth of field was 8.1 nm at a 200 keV acceleration voltage and a 23 mrad probe convergence angle.14 GaN grows in ‘island mode’ during the early stages due to lattice mismatch with the hetero-substrate.10 Microcolumns during the initial growth stage merge with each other and introduce dislocations as their crystal orientations are nearly identical. This GaN can be regarded as a single crystal containing many dislocations. As TEDs (be=aj=⅓, j=1,2,3) parallel to the c-axis exhibit adequate structure for the coalescence of microcolumns, they occupy the majority of dislocations (even up to 80%); TSDs (Burgers vector b=) comprise a small portion of the total, and the remaining are mixed type threading dislocations (bm =⅓).10,15


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Figure 1. Frequent linear arrays consisting of pure edge type dislocations in electron images of cross-sectional WBDF (a) and plan-view STEM (b). TEDs of the linear arrays are identically inclined in the direction normal to the linear array (b), and their Burgers vectors were normal to the linear array of dislocation.

In the cross-sectional WBDF, most of the dislocations were inclined more than 10°. Notably, groups of equally spaced TED arrays appeared frequently. Applying surface relaxation contrast techniques in plan-view, most dislocations exhibited line contrast of TEDs (Figure S1), and many were aligned like the equally spaced TEDs in the WBDF analysis (Figure 1a). Such a linear array of dislocations has been reported in annealed GaN.15 The linear arrays of dislocations were confirmed to consist of only TEDs using atom force microscopy analysis, which is in accordance with our analysis. Such methods of TED


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alignment reduce elastic energy.16 Here, dislocations on the linear array also were utilized to determine the atomic structure of the inclined TED. The dislocation density was 6.8*108/cm2 using plan-view bright field TEM images; the constituent types were analyzed to be 86% edge, 13% mixed, and 1% screw from the surface relaxation contrast images. Typical TEDs are parallel to the c-axis, and their atomic structures do not change with growth. They have been found to possess three types of core structures referred to as 8-, 4-, and 5/7-atomic cores.17,18 Here, the number denotes the number of atoms in the atomic rings of the respective core. The inclination of TED indicates its periodical movement toward the

a-axis by the successive climb. In other words, considering the line of broken bonded atoms, the core lying parallel to the c-axis is cut into small segments and shifted periodically to the

a-axis by successive climbs with the form of stairs. The atomic structure would change from its initial structure because the inclined core arrives on a bulk area with no lattice defect. Microscopic studies require high resolution with respect to the c-axis to distinguish the atomic structure of the individual segments. The applied depth of field of 8.1 nm was similar to that used to investigate inclined mixed type dislocations in GaN.14 To reveal the atomic structure of the core segment, we first observed the plan-view specimen with the Cs-corrected STEM. In the STEM view on the linear array of dislocation, extra-half planes of the constituting TEDs were parallel to the linear array. The Burgers vectors were be =⅓ as in the conventional TEDs; the exact locations of the cores were uncertain. A highly crowded atomic arrangement was observed around the estimated core positions. In total, 57 TEDs on 22 linear dislocation arrays were examined; 47 TEDs exhibited an atomic arrangement similar to that shown in Figure 2a. These arrangements formed rectangles of several nanometers in width and length. Because the cores were shredded to short segments by the successive climbs, the crowded atomic region probably contained cores both prior to and after a climb. As edge dislocations


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are located between the tensile and compressed strain region and GaN is a binary material, the interatomic distances (Ga-N) would change significantly around the core. Thus, upon varying the electron probe focus along the c-axis, observation of the non-overlapped atomic arrangement before/after the climb of the crowded region was expected. We recorded the atomic images by decreasing the focus over 2 nm intervals along the c-axis and distinguished two different (0002) planes of the lower and upper parts of the thin plan-view specimen at the crowded atomic region. (Figure 2a red circles → blue circles) The overlapped atomic region was located near the tensile strain region below the dislocation core, and the displacement of two overlapped (0002) planes was equal to one-half Burgers vector of TED (=½be). The atomic arrangement of the overlapped region can be simply tested by overlapping two TED atomic structure drawings. (Like Figure 3) Upon applying the displacement of ½be, the overlap of two (0002) perfect TED locations does not produce the atomic arrangement demonstrated in the STEM observation. Meanwhile, we can infer a partial TED from the observed displacement (½be) of two (0002) planes, such as in Figure 3. As will be discussed below, a partial dislocation has been introduced to study the inclined mixed-type dislocation. The overlapped (0002) planes of two partial TEDs with a displacement of be produce an atomic structure corresponding to the observed atomic arrangement. From the dislocation contrast shown in Figure 1b, the direction of TED inclination is mostly toward , toward the Burger’s vector. The area of the overlapped atomic region will expand toward proportional to the climb number. We acquired images that show clearly the intersection of two (0002) planes, which have atomic structures of the partial dislocation, similar to the inclined mixed-type dislocation. (Figure S2)


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Figure 2. Overlapping atomic structures appeared at most dislocations comprising linear arrays in the STEM view. From optical sectioning, structures were found to consist of two (0002) planes with a displacement of ½be, which was half the Burgers vector of TED. (a) (Figure S2) The cross-sectional view showed a discrepancy of the atomic arrangement of the identical (0002) plane. Displacement between the compressive and tensile strain regions of TEDs was measured to be ½be. (b) (Figure S3) Here, contrast was attributed to only Ga atoms because the intensity of STEM in Z-contrast was proportional to ~Z1.7 (Z=atomic number). The red and blue circles in the micrographs indicate Ga atoms belonging to the different atomic planes of (0002) for (a) and (11-20) for (b).

Next, we investigated the atomic overlap caused by differences in interatomic Ga-N distances between the tensile and compressive strain regions in the view direction using cross-sectional specimens. That is, we tried to observe the displacement between the two (11-20) planes of the respective strain regions. If both strain regions were located at the respective front (back) and back (front) regions of the thin cross-sectional specimen foil, we


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could discriminate the atomic arrangement of the (11-20) planes between regions using STEM optical sectioning. Here, the displacement also refers to the difference in GaN atomic arrangement between the tensile and compressive strain regions of identical (0002) planes. The method is different from that of plan-view optical sectioning, which showed the displacement between two overlapping (0002) planes of two different core segments. In the STEM observation, the dislocations revealed clear lines at low magnification, but non-overlapped atomic structures at high magnification. Thus, the lines do not appear to correspond to the exact locations of the dislocation core. The overlapped structures were found at the ends of some lines disappearing in the middle of the GaN film, caused by the inclination of the dislocations. The overlapping structure, i.e., the atomic arrangement of two overlapping (11-20) planes, was resolved by varying the electron probe focus. The displacement of the two (11-20) planes was about half the TED Burgers vector (½be). (Figure 2b and Figure S3) This implies that the displacement of the atoms in both the tensile and compressive strain regions from each identical (0002) plane is ½be, identical to the Burgers vector of a partial TED in the plan-view result.

Figure 3. Four successive climbs of the 4/8/4/8/4 partial TED with a displacement of ⅓ (=be) (arrows) produce a similar atomic structure to that shown in the STEM view of Figure 2a. Large circles represent gallium, and small circles are nitrogen atoms.


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Inclined TED has been studied in multiple AlxGa1-xN layers grown on sapphire.19,20 AlxGa1-xN shares the crystal structure and mechanism of dislocation of GaN; thus, it is used for the buffer layer to reduce lattice mismatch with the heterosubstrate during GaN growth.3,4 The substitution of small Al atoms for large Ga atoms induces surface vacancies and compressive stress simultaneously. The degree of inclination is proportional to Al concentration. Follsted et al. suggested the surface vacancies are the cause of the TED inclination, which is referred to as the "surface-mediated climb mechanism".19 This mechanism accompanies the reduction of an extra half-plane of TED. The GaN/Si wafer remains highly convex during growth due to the buffer layer. Thus, the surface area increases with GaN growth. If the increased area is filled with vacancies at the core, the TED climbs according to the surface-mediated climb mechanism. In terms of induced stress, Cantu et al. proposed the “stress relieves climb mechanism”.20 The misfit dislocation segments (jog or kink) created by TED climb release the compressive stress created during growth. Correspondingly, the curvature induced by the buffer layers decreases with growth in GaN/Si.3 As seen in Figure 1b, most of the TEDs composing the linear arrays climb almost normal to the extra half-plane. The present inclination is seen to relieve the stress better than mediate a surface vacancy because the ‘stress relieves climb’ mechanism does not concern the atomic vacancy. The atomic structures of partial TED here were introduced from the inclined mixed type dislocations studied by Hirsch et al.21 The 4/8/4/8/4 structure was first introduced by Xin et al.22 The repeated four and eight atom-rings form a stacking fault (SF) area between the two partial dislocations. Because the screw component of the dislocation is not seen at the view (the Burgers vector of the screw dislocation is parallel to the beam direction), the 4/8/4/8/4 structure could be composed of TEDs or mixed type dislocations. The stability


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of the 4/8/4/8/4 atomic structure was investigated in the inversion domain boundary (IDB) by Northrup et al.23 The energy of the IDB is as high as that of the (-21-10) surface because Ga atoms form bonds with other Ga atoms (Ga-Ga), as in the present partial TED. Hirsch et al. adopted an Eshelby twist for the 4/8/4/8/4 structure to avoid unstable Ga-Ga bonding.21 Eshelby twist is the torsional deformation of a crystal resulting from a screw dislocation along the c-axis14, which can change Ga-Ga bonds to Ga-N bonds in the SF of the dislocation. If so, the 4/8/4/8/4 structure becomes a partial mixed type dislocation with a Burgers vector b=½(a+c)+½(a+c). Yang et al. showed the screw component of partial mixed type dislocations in the cross-sectional STEM observation.24 The observed (0002) planes shifted toward the ±½(a+c) direction, whereas the (0002) plane in our cross-sectional observation shifted only toward the direction of the a-axis. Nonetheless, TED possessing unstable Ga-Ga structures at the view is the current experimental result. We think the reason for the unstable Ga-Ga bonding is related to the complementarity between strain and bonding energy in a dislocation. Unlike the IDB studied by Northrup et al., the dislocation grew under high stress by the highly curved wafer in this study. If a perfect TED unavoidably climbs according to the inclination mechanism, the core arrives on the GaN area with no broken bonds. The extra half-plane of a TED is a wide lattice plane comprised of Ga and N atom pairs, requiring long distance movement for a climb. Thus, dislocation will suffer serious strain around the core if the atomic structure does not change. Dislocation energy is the sum of strain energy and atomic bonding energy. The complementarity between the two energies appears in 8- and 5/7-core TED.25 Both TEDs possess one and two broken-bonded atomic columns toward the c-axis, respectively. The 5/7is a relaxed structure from an 8-core TED, and its Ga atoms in the two broken-bonded atomic columns are facing each other, forming weak Ga-Ga metallic bonds. Nonetheless, the energies of both TEDs are similar.5,25,26 Notably, the 5/7-core TED becomes more stable


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under compressive stress.25 It shows how the TED relaxes the strain energy with the creation of broken bonds. In addition, the GaN grew under high stress caused by the buffer layers. That is, the strain and stress are thought to be responsible for the abnormal Ga-Ga bonding differences from the previous studies. We are further investigating the energy of Ga-Ga bonding with ab-initio calculation. The excellent optoelectronic properties of GaN amid a high density of dislocations (~109/cm2) are attributed to the electronic inactivity of the TEDs. With regard to the other IIIV compound semiconductors such as GaAs, the dislocation density is six orders of magnitude lower than that of GaN for the same luminous efficacy.1 In GaN, TEDs are the dominant dislocation types, occupying up to ~80% of the total dislocations.15 The electronic inactivity of the TEDs is due to its broken-bonded atoms possessing a (10-10) surface-like structure.5,26 The (10-10) GaN surface is a dimerized structure of sp2-p3 (Ga-N) bonds with no defectlevels within the bandgap. However, inclination seems to induce the TEDs to be electrically active from the current result. The Ga-Ga interatomic distance of the present 4/8/4/8/4 partial TEDs was measured to be ~2.0 Å in the STEM view. The metallic Ga bond length was 2.44 Å. Notably, this indicates the generation of electron leakage paths due to strong Ga metallic bonding, which is thought to be a major reason for the inferior luminous efficacy of GaN/Si devices. In general, the atomic structures of dislocations are perceived to be identical to those of GaN grown on sapphire. The electron micrographs in the reports on GaN/Si show highly inclined dislocations.6–9 On the other hand, highly inclined dislocations are observed also within the vicinity of quantum wells where impurities are heavily doped in GaN/sapphire. In summary, we investigated inclined threading edge dislocation, which is the majority of GaN dislocation grown on Si(111). Both plan and cross-sectional electron microscope observation revealed that the dislocation consisted of two partial edge dislocations forming


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the 4/8/4/8/4 structure. The observed structure possessed a Ga-Ga atomic configuration different from the known 4/8/4/8/4 mixed type dislocation constituted of purely Ga-N bonds. This Ga-Ga atomic configuration is considered to form Ga metallic bonds and be one of the important reasons for the inferior efficacy of GaN grown on Si(111) relative to GaN grown on sapphire.

AUTHOR INFORMATION Corresponding Author *[email protected] ORCID Mino Yang: 0000-0002-3860-602X

Notes The authors declare no competing financial interest.

Acknowledgement This research was supported by the Basic Science Research Program through through the National Research Foundation of Korea (NRF-2016R1C1B1013667), and by Korean Basis Science Institute (E38300).

Supporting Information is available free of charge via the Internet at http://pubs.acs.org


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: Surface relaxation contrast image,and plan-/cross-sectional optical sectionning images (docx)

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