ARTICLE pubs.acs.org/JPCC
Strain-Induced ZnO Spinterfaces C. S. Ong,† T. S. Herng,† X. L. Huang,† Y. P. Feng,‡ and J. Ding*,† †
Department of Materials Science and Engineering, National University of Singapore, 7 Engineering Drive 1, Singapore 117574, Singapore ‡ Department of Physics, National University of Singapore, 2 Science Drive 3, Singapore 117542, Singapore
bS Supporting Information ABSTRACT: A series of undoped ZnO films of different thicknesses was grown on different substrates over a range of different temperatures and oxygen partial pressures. Notably, ferromagnetism was detected in very thin ZnO films (∼20 nm), and its magnetic ordering was also found to be thermally stable up to 800 °C. To our surprise, magnetic ordering was destroyed as the ZnO overlayer grew thicker, just as its in-plane compressive strain was released and the ZnO/substrate interface damaged by misfit dislocations. The source of magnetism was found to be due to neither defects in the bulk of the ZnO overlayer nor the bulk of the substrate. Experimental results showed that strain at the ZnO/substrate interface led to a strain-induced magnetic effect. Using firstprinciples ab intio calculation, we confirmed that strain at the ZnO/substrate interface stabilizes zinc vacancy defects, which are magnetic. Ferromagnetic ordering is a result of the coupling of unpaired electron spins originating from the oxygen 2p orbitals surrounding the zinc vacancies.
’ INTRODUCTION Defects-free ZnO is a semiconductor that is widely known to be nonmagnetic.1 Recently, exciting research in the field of diluted magnetic semiconductors (DMS) has unveiled the ability to attain intrinsic ferromagnetism in ZnO at room temperature through defect mediation.2 Bulk ferromagnetism has been reported in ZnO doped with Mn, Cu, Co, C, Li, and N and undoped ZnO.3 10 In many cases, ferromagnetism has been attributed to the coupling of unpaired electron spins that was introduced as a result of crystal defects. On the other hand, concomitant advances in the fields of spinterface science and topological insulators have also stirred a flurry of activities in twodimensional (2D) ferromagnetism and electron gas research.11 16 Magnetic effects at the interface of nonmagnetic oxides as well as ferromagnetism on the surface of hydrogenated-ZnO film have also been reported.17,18 From a broader perspective, these research discoveries bode well for the advent of ground-breaking spintronic semiconductor devices, which endeavor not only to utilize the charge of electrons to process information like traditional electronic devices, but also to take advantage of the spin of electrons, to store information while processing them at the same time. A ferromagnetic semiconductor interface plays an important role in the success of these devices as it improves the efficiency of spin injection across the interface.19 21 In this study, a series of undoped ZnO films of different thicknesses (20 200 nm) was grown on different substrates over a range of different temperatures (room temperature 600 °C) and oxygen partial pressures (6 10 6 to 1 10 3 Torr). Notably, ferromagnetism was detected in very thin ZnO films (∼20 nm) and its magnetic ordering was also found to be thermally r 2011 American Chemical Society
stable up to 800 °C. The magnetic ordering was destroyed as the ZnO overlayer grew thicker, just as its in-plane compressive strain was released and the ZnO/substrate interface damaged by misfit dislocations. Magnetism was found to be confined at the strained ZnO/substrate interface and could neither be attributed to defects in the bulk of the ZnO overlayer nor the bulk of the substrate. Using first-principles ab intio calculation, we found that strain at the ZnO/ substrate interface stabilizes zinc vacancy defects, which are magnetic. Ferromagnetic ordering is a result of the coupling of unpaired electron spins originating from the oxygen 2p orbitals surrounding the zinc vacancies.
’ EXPERIMENTAL RESULTS Undoped ZnO films were prepared by pulsed-laser deposition (PLD). The ZnO target was prepared by sintering ZnO (99.99% purity, Sigma-Aldrich) at 1000 °C for 12 h. A KrF excimer laser operating at 248 nm and a fluence of 1:8 J cm 2 was used to produce the laser plume. Magnetic measurements were made using Superconducting Quantum Interference Device (SQUID; Quantum Design, MPMS, XL-5). External magnetic field was applied in the in-plane direction. Structural characterizations were performed using high-resolution X-ray diffraction (HRXRD; Bruker D8 Discover High-Resolution X-ray Diffraction System), while optical characterizations were made using Raman spectrometry (LabRam HR 800). Both bulk ZnO and bulk substrates were individually determined to be nonmagnetic by SQUID. Received: June 4, 2011 Revised: October 13, 2011 Published: November 30, 2011 610
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direction [100], expanded perpendicularly in the out-of-plane direction [002]. Using the literature value of the (100) peak of a fully relaxed ZnO (2θ = 31.77°), a quick calculation showed that the 200 nm film was under an in-plane compressive strain of 0.5%, whereas the 20 nm film was under a greater in-plane compressive strain of 1.1%. Temperature study was then carried out to study its effect on the magnetic stability of 20 nm films that were prepared at an oxygen partial pressure of 6 10 6 Torr and a substrate temperature of 350 °C. Magnetization of the 20 nm film remained at ∼4 emu/cm3 after 8 h of 800 °C open-air annealing (Figure 3a). In contrast, magnetization of thicker films was always significantly suppressed after annealing (Figure 4a, inset). To have a better understanding, XRD analysis was further performed on the 20 nm ZnO films. The (002) peak was found to be located at the same position before and after annealing, indicating that for 20 nm films, open air-annealing could not relieve the strain (Figure 3a Inset). This was because very thin films are low in local strain and stress densities; strain energy that could be released by the nucleation of misfit dislocations would be smaller than the energy needed to nucleate these dislocations. In fact, the magnetization of 20 nm films was also insensitive to the substrate temperature: Another 20 nm film that was prepared at room temperature possessed an MS of ∼4 emu/cm3 as well (Figure 3a). Magnetism could not have been due to defects in the bulk of the ZnO overlayer, as Raman spectra showed that the presence of defect peaks at 580 cm 1 (Figure 3b), which corresponded to bulk defects in the ZnO overlayer (Figure 4b,d), was independent of its magnetization. The results indicated that magnetic ordering for 20 nm films was thermally stable and that ferromagnetism was associated with the ZnO/substrate interface. The fact that when ZnO film was grown beyond a certain critical thickness, much of the magnetic ordering was destroyed just as its pent-up in-plane compressive strain was released, pointed to an interface-related magnetic effect. ZnO first grew pseudomorphically, elastically strained to match the interatomic spacing of the substrate surface. As the epitaxial ZnO overlayers grew thicker, strain energy accumulated. Beyond the critical thickness, pent-up strain energy was released, overcoming the energy barrier associated with dislocation nucleation and propagation, leading to misfit dislocation arrays forming at the ZnO/ substrate interface and destroying the magnetic interface for thicker films (Figure 5). As evident from the XRD results, lattice parameters of thicker films were much closer to its free lattice constants (Figure 2a,b) due to film relaxation. To confirm the strain-induced spinterface effect, a series of 20 nm ZnO films was deposited on five different substrates (X-cut quartz, Y-cut quartz, Z-cut quartz, glass, and silicon (001)) at an oxygen partial pressure of 6 10 6 Torr and a substrate temperature of 350 °C. The MS of the ZnO films decreased when they were grown on substrates in the following order: X-cut quartz, Y-cut quartz, and glass (Figure 6). No magnetic ordering was observed when ZnO films were grown on Z-cut quartz and silicon (001). XRD analysis also showed that the (002) peak shifted left (Figure 7a) and the (100) peak shifted right (Figure 7b) as magnetization increased. The results were consistent with our studies above, underlining the importance of strain to induce the observed interfacial ferromagnetism. In particular, SiO2 substrates in the form of X-cut quartz, Y-cut quartz, Z-cut quartz, and glass all led to different and distinct levels of magnetic ordering in ZnO. Notice that all except glass have identical bulk structures, differing only at their surfaces. Since the ZnO overlayers were all equally
Figure 1. Saturation magnetization (MS) of ZnO films as a function of thickness. Colored markers outline the thickness trend for films grown at an oxygen partial pressure of 6 10 6 Torr, while light gray markers outline the trend when films were grown at a much higher oxygen partial pressure of 1 10 6 Torr. Inset contains the magnetic hysteresis loops of ZnO films at different thicknesses when grown at an oxygen partial pressure of 6 10 6 Torr after subtracting the linear diamagnetic substrate contribution.
The ZnO films were also confirmed to be highly textured in the [001] direction using XRD. First, we began with an investigation of the effect of thickness on the magnetic properties of ZnO films. Epitaxial ZnO films were grown on X-cut quartz substrates at an oxygen partial pressure (PO) of 6 10 6 Torr and a substrate temperature of 350 °C. Films of five different thicknesses were prepared: 20, 50, 100, 150, and 200 nm. It was discovered that magnetization was weaker as the films were grown thicker. In Figure 1, we see that thicker films had lower saturation magnetization (MS), with the 20 nm film being the most magnetic and the 200 nm film not being magnetic at all. In addition to the thickness effect, it was noticed that the magnetic properties of undoped ZnO were also sensitive to the oxygen partial pressure (PO). As shown in Figure 1, when prepared at a higher oxygen partial pressure of PO = 1 10 3 Torr, the film was nonmagnetic down to a thickness of 50 nm, which was also in good agreement with our previous work.5 However, when its thickness was reduced to 20 nm, magnetic ordering was detected in the ZnO film. The results suggest that magnetism in ZnO films was related to the film being in a very thin state. From here on, since ferromagnetism is central to this study, our subsequent discussions will focus on ZnO films prepared at PO = 6 10 6 Torr. Out-of-plane XRD measurements of our ZnO films revealed that the predominant (002) peaks of thinner films were always located to the left of thicker ZnO films (Figure 2a). In-plane XRD measurements of the (100) planes, which are perpendicular to the (002) planes, were made by rotating the substrates such that incident X-ray now entered the ZnO epitaxial overlayer in a direction parallel to the ZnO/substrate interface. In contrast to behavior of the (002) peaks, the (100) peaks always shifted right for thinner samples (Figure 2b). Physically, this meant that for thinner samples, there was greater compressive strain along the a- and b-axes, coupled with a lattice expansion along the c-axis, in accordance to Bragg’s law. In fact, this was also a manifestation of the Poisson effect: ZnO when compressed in the in-plane 611
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Figure 2. (a) ZnO (002) XRD peak shifted left for thinner ZnO films. (b) ZnO (100) XRD peak shifted right for thinner ZnO films. Literature values of the (002) and (100) peak positions of a fully relaxed ZnO crystal are at 34.42° and 31.77°, respectively. Photoluminescence and UV visible spectra of the 20 and 200 nm films are displayed in panels c and d, respectively. (e) The transport properties of 20 and 200 nm films. (f,g,h) Scanning electron microscopy (SEM) image, atomic force microscopy (AFM) image, and energy-dispersive X-ray (EDX) spectrum of a typical ZnO overlayer, respectively.
thin, it would have to be the 2D surface periodicity of the substrates at the interface, and not the elemental composition of the substrates or the film, that had played the deciding role in its magnetic ordering. Our experimental studies demonstrated that there was a close correlation between the strain at the ZnO/ substrate interface and its magnetization.
(DFT) as implemented in the Vienna ab initio Simulation Package (VASP) code, using projector augmented wave (PAW) potentials for electron ion interaction and generalized gradient approximation (GGA-PW91) for electron exchange and correlation.22 25 Cutoff energy was set at 400 eV for the plane-wave basis. In all calculations, self-consistency was achieved with accuracy of at least 10 3 and 10 4 eV for the ionic and electronic loops, respectively. Native bulk defects were modeled with periodic supercells consisting of 3 3 2 wurtzite ZnO unit cells, and k points were generated with a 3 3 2 Γ-centered grid based on the Monkhorst-Pack scheme.26 The interfaces were modeled using periodic supercells consisting of 2 8 1 wurtzite ZnO unit
’ THEORETICAL RESULTS To have a better understanding of the origin of ferromagnetism in the ZnO thin film, we carried out first-principles calculation. The calculation was based on the density functional theory 612
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Figure 3. (a) Magnetization curves of 20 nm (very thin) ZnO films after correcting for the diamagnetic signal from the substrate. Magnetization remained unchanged after open-air annealing. It was also insensitive to the substrate temperature. Inset shows the ZnO (002) XRD peak of the film before and after open-air annealing. (b) Raman spectra of very thin ZnO films prepared under different conditions.
Figure 4. ZnO (002) XRD peak for the 150 nm films prepared at different (a) oxygen partial pressures and (c) substrate temperatures. The saturation magnetizations as functions of oxygen partial pressure and substrate temperature are shown in the insets of panels a and c, respectively. ZnO films with higher magnetization were accompanied by left shifts of the (002) peak. Panels b and d respectively show the room-temperature Raman Spectra of ZnO films prepared at different oxygen partial pressures and substrate temperatures. The data revealed the presence of a defect peak at 580 cm 1, which was stronger for films grown at lower oxygen partial pressure and lower substrate temperature. 613
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Figure 5. A schematic drawing of ZnO film growth. Our experimental results showed that a higher degree of magnetic ordering invariably corresponded to a ZnO overlayer under larger in-plane compressive strain, coupled with an out-of-plane relaxation. Film growth dynamics suggest that beyond certain critical thickness, the onset of misfit dislocation nucleation and propagation destroyed the pseudomorphic, strained, and magnetic interface (Figure 4a), disrupting its interfacial magnetic ordering (Figure 4b).
energy of the corresponding perfect supercell, ni, the number of atoms removed, and μi is the chemical potential of the corresponding atom. The actual chemical potential μi depends on experimental growth conditions and ranges from Zn-rich conditions to O-rich conditions. In extreme O-rich condition, μO is given by the energy of O in an O2 molecule, while in extreme Zn-rich conditions, μZn is given by the energy of Zn in bulk hexagonal close-packed (hcp) zinc. In this study, μO in O-rich condition was used as the reference and was set as the upper bound of μO at 0 eV. In addition, for Zn and O particle reservoirs to be at thermodynamic equilibrium with bulk ZnO, the following condition has to be satisfied as well: μZn + μO = EZnO, where EZnO is the energy of ZnO. This set the lower bound of μO to be at 3.04 eV, which is the heat of formation of ZnO. Our calculated lattice parameters and strain agreed very well with our experimental data. For pure unstrained ZnO, the computed lattice parameters of a = b = 3.29 Å, c = 5.29 Å compared favorably with our experimental results and the literature values27 of a = b = 3.25 Å, c = 5.21 Å. Upon the introduction of strain, the calculated lattice parameters agreed well with our experimental data as well(Table 1). Even the signs of the strain in our theoretical model were similar to those observed experimentally (Table 1), confirming the existence of compressive strain along the a- and b-axes and a lattice expansion along the c-axis. Table 2 shows that the experimental trend of increasing strain, when grown on substrates in the following order: X-cut quartz, Y-cut quartz, Si (001) and Z-cut quartz, agreed also with the trend of increasing theoretical maximum strain. In our magnetism study, bulk ZnO and SiO2 were confirmed to be nonmagnetic, as it was also observed experimentally. Oxygen vacancy (VO) is nonmagnetic in both unstrained and strained ZnO (Figure 9). Zinc vacancy (VZn), on the other hand, is magnetic. When a saturating external magnetic field was applied in the in-plane direction, unpaired electron spins mostly originating from the unfilled oxygen 2p orbitals surrounding VZn aligned in the direction of the field and contributed to a magnetic dipole moment of 0.3 μB per defect in both unstrained and strained ZnO. This was because the presence of VZn introduces “dangling holes”. For ZnO, the O 2p states are located more shallowly below the valence band maximum (VBM) than the
Figure 6. Saturation magnetization (MS) of 20 nm ZnO films as deposited on the X-cut, Y-cut, Z-cut quartz, glass, and Si (001) substrates. Their magnetization curves after subtracting the linear diamagnetic contribution from the substrate are shown in the inset.
)
)
cells built on top of 1 3 1 (110)-terminated α-quartz substrate units with an epitaxial orientation relationship of (002)ZnO (110)SiO2 and an in-plane alignment of [100] ZnO [110]SiO2 (Figure 8). The entire supercell slab was 10 Å-thick, separated by a vacuum space of at least 15 Å. Each slab was made up of 4 atomic layers of ZnO that was 5 Å-thick, built on top of a (110)-terminated α-quartz substrate which was also 5 Å-thick. Both surfaces were terminated using pseudohydrogen to saturate the bonds and simulate bulk. The models were tested to have converged with regard to both slab thickness and vacuum gap thickness. k points were generated with a 1 3 1 Γ-centered grid. All atoms in the models were fully relaxed. In our calculation, the formation energy of a defect i in neutral state was calculated by using Ef = Edtot E0tot + Σiniμi, where Edtot is the total energy of the supercell with defects, E0tot is the total 614
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)
)
Figure 7. XRD spectra of 20 nm ZnO films deposited at 350 °C on different substrates. The (002) peaks are shown in panel a, while the (100) peaks are shown in panel b. The trend is consistent with that in Figure 1, i.e., ZnO films with a higher magnetic ordering were always accompanied by left shifts of the (002) peak and right shifts of the (100) peak.
Figure 8. Epitaxial orientation relationship at the ZnO SiO2 interface. (002)ZnO (110)SiO2; [100]ZnO [110]SiO2.
contributing a magnetic dipole moment of 0.3 μB per defect (Figure 9). The spin-density is localized at the ZnO/substrate interface, due to unpaired electron spins from the O 2p orbitals at the interface (Figure 10). A more highly strained interface stabilizes a higher concentration of VZn, leading to a larger number of magnetic dipoles orienting in the direction of the magnetic field when an external saturating magnetic field is applied. This accounts for the higher saturation magnetization (MS) that was experimentally observed for a more highly strained ZnO film (Table 2, Figures 1, 7). Our theoretical calculation showed that the ferromagnetic state has a much lower energy level than the antiferromagnetic state. Different antiferromagnetic configurations were modeled, and they were always found to be unstable and energetically unfavorable. The block Davidson scheme28,29 that was used to search for the electronic ground state always converged to a ground
Zn 3d states. Hence, O 2p orbitals have the tendency to lose electrons more easily than Zn 3d orbitals, leading to spin polarization of the O 2p orbitals. However, in our experiment, the oxygen partial pressure during film growth was very low (PO = 6 10 6 Torr) and VZn could not have formed readily (Ef ∼ 5 eV). First-principles calculation also showed that epitaxial strain in bulk ZnO alone is unable to lower the formation energy of VZn effectively (ΔEf = 0.07 eV) (Figure 9). Next, we considered a strain-induced interfacial magnetic effect. Many different defects-free perfect ZnO/quartz interfaces were modeled. All were found to be nonmagnetic. More information regarding these interface models can be found in the Supporting Information. A single VZn was then introduced at the interface (Figure 10). Interestingly, the formation energy of VZn at the interface is substantially lower than its formation energy in bulk ZnO (ΔEf = 1.3 eV) (Figure 9). It is magnetic too, 615
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Table 1. Comparison between the Experimental and Calculated Lattice Parameters of Unstrained and Strained ZnO Structures computational results experimental results fully relaxed ZnO (Unstrained)
without defect
a (Å)
3.25
3.29
b (Å)
3.25
3.29
with VO
with VZn N.A.
c (Å)
5.21
5.29
c/a
1.60
1.61
volume (Å3)
47.6
49.5
a (Å)
3.22
3.27
3.27
3.27
b (Å)
3.22
3.16
3.16
3.16
c (Å) c/a
5.25 1.63
5.40 1.65
5.40 1.65
5.36 1.64
volume (Å3)
epitaxially strained ZnO as deposited on X-cut quartz (Strained)
percentage change (%)
47.0
48.3
48.3
48.0
a
1
1
1
1
b
1
4
4
4
c
1
2
2
1
c/a
2
3
3
2
volume
1
3
3
3
Table 2. Experimentally Measured In-Plane Strain in ZnO Films as Compared to Their Theoretical Maximaa observed
experimental
theoretical maximum
magnetism?
in-plane strain (%)
in-plane strain (%)
X-cut quartz
yes
1.03
4.11
Y-cut quartz glass
yes yes
1.00 0.70
4.11 N.A.
silicon (001)
no
0.52
1.29
Z-cut quartz
no
0.52
0.81
Figure 10. An atomistic representation of the model illustrated in Figure 5a, obtained using first-principles calculation. The yellow shaded regions mark the spin densities of the electrons. We see that magnetization was mostly due to the ferromagnetic coupling of unpaired electron spins occupying the oxygen 2p orbitals as a result of VZn.33
a
Also shown is the correlation between in-plane strain and the observed magnetism.
antiferromagnetic states, which is small for our case. When another VZn defect was introduced half a supercell length away from the first VZn defect in the ZnO [100]-direction, the interface remained magnetic, further confirming that coupling between the magnetic VZn defects at the interface is ferromagnetic and not antiferromagnetic. Since strain stabilizes a higher concentration of magnetic spin dipoles at the interface, a more highly strained interface will also exhibit a higher residual magnetization when the external magnetic field is removed, as was also experimentally verified (Figure 1). This is because the magnetic remanence is a result of net spin dipoles lying in the direction of easy magnetization in the absence of an external field. Meanwhile, coercive field is a measure of the stability of the residual magnetization. In our samples, the origin of magnetic coupling is consistent, attributed to the coupling of unpaired electron spins from the O 2p orbitals at a 2D ZnO/substrate interface. As a result, the stability of the residual magnetization and the coercive field do not differ significantly and qualitatively from one sample to another, which was also our experimental observation (Figure 1). In fact, defectmediated magnetic coupling is typically characterized by a small coercive field. The shapes of our magnetization curves and the magnitudes of the magnetization are also highly characteristic of defects-mediated magnetism.8,30 32
Figure 9. Formation energies of different defects as functions of the oxygen chemical potential μO. Actual μO depends on the experimental growth conditions, which range from O-rich to Zn-rich conditions. Indicated on the same plot are also the magnetic dipole moments of the defects.
state that is ferromagnetic. At temperatures above 0 K, magnetic behaviors can be determined by the ratio of thermal energy in relation to the energy difference between the ferromagnetic and 616
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’ CONCLUSION In conclusion, we demonstrated that the amount of strain at the ZnO/substrate interface significantly affected its interfacial magnetic ordering. Very thin films were always magnetic, and magnetic ordering at the interface was thermally stable up to 800 °C. When equally thin ZnO films were grown on substrates that were differently terminated but of the same chemical elemental composition, magnetizations were distinctly different. A larger saturation magnetization was always observed when the in-plane compressive strain was greater. This indicated that the observed magnetism was due to neither the bulk of ZnO nor the bulk of the substrate, but was instead a property of the strained ZnO/substrate interface. Using first-principles ab intio calculation, we found that although VZn is magnetic, it would not have formed easily in the bulk of the ZnO overlayer under our oxygenpoor experimental conditions. The presence of compressive inplane strain at the interface substantially lowered the formation energy of VZn, leading to spin densities localized at the interface. Spin polarization was due to unpaired electron spins from the oxygen 2p orbitals. In the presence of an external saturating magnetic field, a more highly strained interface had a higher concentration of unpaired electron spins oriented in the direction of the magnetic field, contributing to the higher saturation magnetization. This magnetism is confined at the ZnO/substrate interface, and is known as the “Spinterface” effect.
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’ ASSOCIATED CONTENT
bS
Supporting Information. Description of the nonmagnetic perfect interfaces models. This material is available free of charge via the Internet at http://pubs.acs.org.
’ AUTHOR INFORMATION Corresponding Author
*Electronic mail:
[email protected]. Tel: (+65) 6516 4317. Fax: (+65) 6776 3604.
’ ACKNOWLEDGMENT This work is supported by the National Research Foundation of Singapore under Grant No. NRF-G-CRP2007-05. ’ REFERENCES (1) Venkatesan, M.; Fitzgerald, C. B.; Coey, J. M. D. Nature 2004, 430, 630. (2) Dietl, T.; Ohno, H.; Matsukura, F.; Cibert, J.; Ferrand, D. Science 2000, 287, 1019. (3) Fukumura, T.; Jin, Z. W.; Kawasaki, M.; Shono, T.; Hasegawa, T.; Koshihara, S.; Koinuma, H. Appl. Phys. Lett. 2001, 78, 958. (4) Ueda, K.; Tabata, H.; Kawai, T. Appl. Phys. Lett. 2001, 79, 988. (5) Herng, T. S.; Qi, D. C.; Berlijn, T.; Yi, J. B.; Yang, K. S.; Dai, Y.; Feng, Y. P.; Santoso, I.; Sanchez-Hanke, C.; Gao, X. Y.; Wee, A. T. S.; Ku, W.; Ding, J.; Rusydi, A. Phys. Rev. Lett. 2010, 105. (6) Lee, H. J.; Jeong, S. Y.; Cho, C. R.; Park, C. H. Appl. Phys. Lett. 2002, 81, 4020. (7) Pan, H.; Yi, J. B.; Shen, L.; Wu, R. Q.; Yang, J. H.; Lin, J. Y.; Feng, Y. P.; Ding, J.; Van, L. H.; Yin, J. H. Phys. Rev. Lett. 2007, 99, 127201. (8) Yi, J. B.; Lim, C. C.; Xing, G. Z.; Fan, H. M.; Van, L. H.; Huang, S. L.; Yang, K. S.; Huang, X. L.; Qin, X. B.; Wang, B. Y.; Wu, T.; Wang, L.; Zhang, H. T.; Gao, X. Y.; Liu, T.; Wee, A. T. S.; Feng, Y. P.; Ding, J. Phys. Rev. Lett. 2010, 104, 137201. 617
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