Surface Chemistry and Interface Formation during the Atomic Layer

Article pubs.acs.org/cm

Surface Chemistry and Interface Formation during the Atomic Layer Deposition of Alumina from Trimethylaluminum and Water on Indium Phosphide Christoph Adelmann,*,† Daniel Cuypers,†,‡ Massimo Tallarida,§ Leonard N. J. Rodriguez,† Astrid De Clercq,†,∥ Daniel Friedrich,§ Thierry Conard,† Annelies Delabie,†,‡ Jin Won Seo,∥ Jean-Pierre Locquet,⊥ Stefan De Gendt,†,‡ Dieter Schmeisser,§ Sven Van Elshocht,† and Matty Caymax† †

Imec, Kapeldreef 75, B-3001 Leuven, Belgium Departement Chemie, Katholieke Universiteit Leuven, Celestijnenlaan 200F, B-3001 Leuven, Belgium § Brandenburgische Technische Universität, Angewandte PhysikSensorik, Konrad-Wachsmann-Allee 17, D-03046 Cottbus, Germany ∥ Departement Metaalkunde en Toegepaste Materiaalkunde, Katholieke Universiteit Leuven, Kasteelpark Arenberg 44, B-3001 Leuven, Belgium ⊥ Departement Natuurkunde en Sterrenkunde, Katholieke Universiteit Leuven, Celestijnenlaan 200D, B-3001 Leuven, Belgium ‡

ABSTRACT: The surface chemistry and the interface formation during the initial stages of the atomic layer deposition (ALD) of Al2O3 from trimethylaluminum (TMA) and H2O on InP(100) were studied by synchrotron radiation photoemission spectroscopy and scanning tunneling microscopy. The effect of the ex situ surface cleaning by either H2SO4 or (NH4)2S was examined. It is shown that the native oxide on the InP surface consisted mainly of indium hydrogen phosphates with a P enrichment at the interface with InP. After a (NH4)2S treatment, S was present on the surface as a sulfide in both surface and subsurface sites. Exposure to TMA led to the formation of a thin AlPO4 layer, irrespective of the surface cleaning. The surface Fermi level of ptype InP was found to be pinned close to midgap after H2SO4 cleaning and moved only slightly further toward the conduction band edge upon TMA exposure, indicating that the AlPO4/InP interface was rather defective. (NH4)2S passivation led to a Fermi level position of p-type InP close to the conduction band edge. Hence, the InP surface was weakly inverted, which can be attributed to surface doping by S donors. TMA exposure was found to remove surface S, which was accompanied by a shift of the Fermi level to midgap, consistent with the removal of (part of) the S donors in combination with a defective AlPO4/InP interface. Further TMA/H2O ALD did not lead to any detectable changes of the AlPO4/InP interface and suggested simple overgrowth with Al2O3. KEYWORDS: atomic layer deposition, InP, Al2O3, trimethylaluminum, interface formation

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gap,4−8 which cannot be passivated by forming gas annealing as can their Si−SiO2 counterparts. Hence, III−V semiconductor based MOSFETs will require a fully deposited gate dielectric without the presence of a native interfacial oxide. Such an interface can be obtained by in situ deposition of the gate dielectric onto clean III−V surfaces. Low-defect interfaces have been demonstrated by in situ molecular-beam-deposited Ga 2 O 3 9 and Gd 2 O 3 10 on GaAs(100); however, such processes have still to be demonstrated on large-scale wafers. For practical manufacturable processes, an ex situ gate dielectric deposition is strongly favored. Because of the rapid oxidation of III−V semiconductors in air, any dielectric

INTRODUCTION The integration of III−V semiconductors, such as GaAs, InGaAs, or InP, into metal−oxide−semiconductor field-effect transistors (MOSFETs) has recently gained renewed interest since the scaling of the conventional silicon-based technology becomes increasingly limited by silicon’s intrinsic properties.1 The main interest in III−V semiconductors lies in the fact that they offer much higher bulk carrier mobilities than silicon. This can lead to an increased MOSFET drive current and thus to an improved performance.1−3 A key obstacle to the manufacturing of high-performance MOSFETs using III−V semiconductors is the long-known observation that their native oxides are generally poor dielectrics, unlike the case of Si. In addition, the interface between III−V semiconductors and their native oxides is highly defective with a large density of interface states within the band © 2013 American Chemical Society

Received: December 20, 2012 Published: February 18, 2013 1078

dx.doi.org/10.1021/cm304070h | Chem. Mater. 2013, 25, 1078−1091

Chemistry of Materials

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in 1.8 M H2SO4 for 5 min or passivated in 2.9 M (NH4)2S for 5 min, followed in each case by a 3 min rinse in deionized H2O. The samples were then blow-dried and transferred to the ALD and measurement apparatus in an inert N2-purged environment. The total air exposure was minimized to a few minutes only. The ALD of Al2O3 from TMA and H2O was carried out in situ in a custom-built ALD chamber, developed at the BTU Cottbus, with a base pressure of 10−9 mbar34 at 300 °C. SRPES measurements were carried out in situ at the U49-2/PGM-2 beamline at the BESSY II synchrotron radiation facility within the Helmholtz-Zentrum Berlin. The base pressure of the photoemission chamber was on the order of 5 × 10−10 mbar. The X-ray photons were monochromatized by a planar grating monochromator in an energy range between 150 and 690 eV. The resolution of the monochromator was ΔE/E ≈ 104. The photoelectrons were detected using a Specs Phoibos 150 analyzer with a pass energy of 10 eV at an emission angle of 45°. All peak fitting was performed within the Thermo Avantage software using pseudo-Voigt (Lorentzian−Gaussian sum) functions. Lorentzian broadenings were on the order of 100 meV; Gaussian broadenings were then allowed to vary to account for the experimental peak widths. Initial values for peak broadenings and in particular for the spin−orbit splittings were obtained from the literature.35 The branching ratios were set to the theoretical values. Within a series of spectra (after different surface preparations and/or for different photon energies), a consistent set of components was determined. Within such a set, chemical shifts (CSs) and peak broadenings were allowed to vary within a 30 meV interval, while spin−orbit splittings were fixed. All spectra were then fitted by only varying the intensities of the components. When in some cases poor or inconsistent fits were obtained, the CSs were also allowed to vary, but again the peak broadenings were kept within tight boundaries of 30 meV. This is discussed in the text. Scanning tunneling microscopy was performed in an Omicron large sample SPM system with a base pressure of 5 × 10−11 mbar. The depicted images were acquired at a bias of 2 V with currents in the 100 pA range. For the as-cleaned samples, the sample preparation procedure was identical to that used for the SRPES measurements. To study the effect of TMA on the InP surface morphology, H2SO4cleaned InP was exposed to TMA at 250 °C in a home-built ultra-highvacuum (UHV)-compatible ALD reactor with a base pressure of 1 × 10−8 mbar and subsequently transported in a vacuum suitcase to the STM chamber without intermittent air exposure. Additional Al2O3 layers were grown from TMA and H2O on InP(100) in an ASM Pulsar ALD reactor at 300 °C after surface preparation identical to that for the SRPES measurements. On these samples, transmission electron microscopy (TEM) was carried out in a Tecnai F30 electron microscope operating at 200 kV in imaging mode and 300 kV in high-angle annular dark field (HAADF) mode. TEM specimens were prepared by focused ion beam sputtering. Additionally, time-of-flight secondary ion mass spectrometry (ToF-SIMS) was performed using an IONTOF-IV instrument with a Ga+ source operating at a beam energy of 15 keV (area of 80 × 80 μm2) for analysis and a Xe+ beam with an energy of 350 eV (sputter area of 400 × 400 μm2) for depth profiling. To obtain ALD growth curves, the deposited Al area density was determined by inductively coupled plasma mass spectrometry (ICPMS) using an Agilent 7500cs system. The deposited Al2O3 was dissolved by etching the wafers in 20 mL of a 1.2 M HCl/0.03 M HF mixture for 10 min. The etched sample area was delimitated by a Kalrez O-ring with a calibrated area of 14.52 cm2. The 27Al concentration in the solution was then analyzed by ICP-MS. The presence of 115In in the solution indicated that the Al2O3 layer was completely etched down to the substrate. To obtain absolute area densities, the measured 27Al concentrations were calibrated against identical measurements on ALD of Al2O3 on Si substrates assuming a steady-state growth per cycle (GPC) of 4 Al atoms nm−2.36

deposition process onto an air-exposed III−V semiconductor surface will have to remove the native oxide in situ. Recently, the observation of the removal of the native oxide on III−V arsenides such as GaAs or InGaAs during the atomic layer deposition (ALD) of dielectrics such as Al2O3 (using trimethylaluminum (TMA) and H2O)11−17 or HfO2 (using tetrakis(ethylmethylamino)hafnium (TEMA-Hf) or HfCl4 in combination with H2O)13,18,19 has led to the exciting prospect of native oxide free interfaces between the dielectric and the III−V semiconductor. By contrast, it is less clear whether the strong reduction of the interfacial arsenic and gallium oxides is concurrent with a reduction in the interface state density. While an unpinning of the Fermi level has been observed by scanning tunneling spectroscopy,20 other measurements using X-ray photoelectron spectroscopy have found little variation of the pinned Fermi level upon exposure of GaAs surfaces to TMA.17,21,22 However, it should be noted that no universal relation between Fermi level position and surface or interface defect density exists. In particular, the substrate doping density is of crucial importance and renders the direct comparison of different studies difficult.23,24 Thus, the classification of surfaces or interfaces as “pinned” or “not pinned” may not be unequivocal. The published data are all compatible with an interface defect density (at least) in the low 1012 cm−2 range after the in situ removal of the native oxide. For III−V phosphides, in particular for InP, the situation is qualitatively different because of the different oxidation thermodynamics.25,26 In contrast with III−V arsenides, thermodynamically stable indium phosphates with good dielectric properties have been obtained on InP substrates.27−31 However, the interface defect density of such phosphate/InP interfaces was still much higher than desirable for advanced metal−oxide−semiconductor (MOS) devices.32 Therefore, the removal of the native oxide (phosphate) in situ before the deposition of the dielectric may be considered necessary for InP also. The reaction between oxidized InP surfaces and TMA has recently been studied by in situ X-ray photoelectron spectroscopy (XPS).33 It was found that the native oxide on InP was not removed by TMA but transformed into a different component, which was attributed to P2O5. In this paper, we comprehensively study by synchrotron radiation photoemission spectroscopy (SRPES) and scanning tunneling microscopy (STM) the surface chemistry and structure as well as the interface formation during the first steps of the ALD of Al2O3 using TMA and H2O. Both the effects of ex situ cleaning by H2SO4 and ex situ passivation by aqueous (NH4)2S on the initial InP surface and the subsequent ALD were considered. We demonstrate that the native oxide reacts with TMA by forming a thin AlPO4 layer. This layer was stable under subsequent H2O pulsing and further ALD. While the presence of such a thin wide band gap oxide in the gate stack is not necessarily detrimental to MOS devices, we also demonstrate that the interface quality between InP and AlPO4 was not improved with respect to the interface between InP and its native oxide. Hence, the suppression of the AlPO4 formation, e.g., by optimized S-based passivation schemes, may be necessary to obtain high-quality MOS devices on InP.

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EXPERIMENTAL DETAILS

All experiments were carried out on Zn-doped p-type (p ≈ 3 × 1017 cm−3) InP(100). Prior to ALD, the samples were either cleaned ex situ 1079



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Figure 1. P 2p SRPES spectra (a−c) using a photon energy of hν = 250 eV and (e−g) using a photon energy of hν = 660 eV. (a) and (e) show spectra after H2SO4 cleaning and degassing at 300 °C, whereas (b) (c), (f), and (g) show spectra after exposure to TMA. The spectra were fitted with the seven components described in the text. (d) and (h) show the evolution of the relative spectral weight of the components as a function of the number of TMA pulses at photon energies of hν = 250 and 660 eV, respectively.

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The P5+ CS (4.20 eV) found in our experiments was significantly lower than the CS of 5.0−5.9 eV reported for bulk InPO4 or In(PO3)3 with respect to bulk InP.40−42 It should however be noted that, for such ultrathin dielectric films (for the quantification of the oxide thickness, see below) on semiconducting substrates, the CSs can deviate strongly from bulk values. This was studied in detail for SiO2/Si43 and attributed to the effect of the dielectric discontinuity at the interface.44 However, arguments against the presence of a true surface phosphate on InP after H2SO4 cleaning can be inferred from the valence band spectra in Figure 2a. The presence of

RESULTS AND DISCUSSION The paper is organized as follows: we first describe the composition and morphology of the initial InP surface after H2SO4 cleaning, followed by the effect of TMA exposure and the chemistry of interface formation. Subsequently, the surface composition and morphology of InP after (NH4)2S passivation is discussed, again followed by the effects of TMA exposure and the chemistry of interface formation. We finally discuss the effects of a subsequent H2O pulse as well as those of continued ALD. TMA Exposure of H2SO4-Cleaned InP. InP Surface after H2SO4 Cleaning and Degassing. The P 2p SRPES spectra of InP after ex situ H2SO4 cleaning and degassing at 300 °C are shown in Figure 1a,e for photon energies of hν = 250 and 660 eV, respectively. The P 2p spectrum was mainly characterized by the P−In substrate peak and a high-valence component (P5+) with a CS of the binding energy (BE) of 4.20 eV with respect to the binding energy of the InP substrate. In addition, components with weaker spectral weight were visible in the spectra, which can be attributed to elemental P0 (CS = 1.25 eV) and a P(2±Δ)+ suboxide (CS = 2.80 eV). The approximate suboxide oxidation state of 2+ was deduced assuming a linear dependence of the chemical shift on the oxidation state. The consistent modeling of the P 2p spectra for both hν = 250 eV and hν = 660 eV required the introduction of two surface components (labeled S1 and S2 with CSs of 0.29 and −0.41 eV, respectively) to obtain good fitting results with a common set of peak parameters (including peak widths and shapes). In the literature, the P 2p photoemission spectra of native surface oxides on InP typically show P5+ components with CSs in the range of ∼4−6 eV, rather similar to our data. The interpretation of these components is somewhat ambiguous, owing to the complexity of the phosphate chemistry. In thermodynamic equilibrium, one would expect the surface oxide to consist of the orthophosphate InPO4.25,26 Indeed, the formation of InPO4 films has been observed after thermal oxidation of InP at high temperature.4,25,37−39

Figure 2. (a) Valence bands and (b) a close-up of the valence band edges of InP after H2SO4 cleaning and degassing as well as after exposure to TMA. In (a), the data are offset for clarity. The photon energy was hν = 150 eV.

phosphates has been shown to lead to the appearance of a characteristic peak around 15 eV.45−47 For phosphates, this peak appears at higher BEs than observed for oxides.48 The valence band spectrum clearly shows that such a peak was absent after H2SO4 cleaning and degassing of the InP surface. It is clear that the native surface oxide of InP after H2SO4 cleaning does not necessarily correspond to well-defined and stoichiometric phosphates, as remarked before.40 Furthermore, the effect of the aqueous environment during the H2SO4 cleaning may not be negligible because of the strong hygroscopy of many phosphates. Indeed P5+ components with even lower CSs of 3.9 eV have been observed after 1080



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Figure 3. (a−c) In 3d5/2 and (e−g) In 4d SRPES spectra (photon energy hν = 660 eV) after H2SO4 cleaning and degassing at 300 °C as well as after exposure to TMA, as specified in the figure. All spectra were fitted with the three components described in the text. (d) and (h) show the evolution of the relative spectral weight of the components as well as the chemical shift of the In3+ component with respect to the In−P substrate component as a function of the number of TMA pulses.

the photoelectrons (∼4 and ∼11 Å for kinetic photoelectron energies of ∼120 and 530 eV, respectively, in InP for the SRPES geometry used),53 the quantitative analysis of SRPES spectra requires a rather accurate knowledge of the depth distribution of the different spectral components. While a detailed quantitative chemical profile of the surface layers is beyond the scope of the paper, some insight can be gained from comparing the spectral weight of the components in spectra of different photoelectron energies. Because the photoelectron inelastic mean free path increases with increasing kinetic photoelectron energy (and thus also for increasing photon energy), a larger relative spectral weight at low photon (low kinetic photoelectron) energies means that the component is closer to the surface. Relative depth plots of the different components in P 2p are shown in Figure 4. The relative depth was defined as the ratio

aqueous surface treatments and attributed to Inx(HPO4)y.49 This thus suggests that the InP surface after H2SO4 cleaning consisted (partially) of indium hydrogen phosphates. Further insight can be gained from the In 3d5/2 and In 4d spectra of InP after ex situ cleaning in H2SO4 and degassing at 300 °C (Figure 3) All spectra could be well accounted for by three components, i.e., the In−P substrate component, In3+ (In 3d5/2 CS = 0.40 eV; In 4d CS = 0.52 eV), and a small amount of In0 (In 3d5/2 CS = −0.37 eV; In 4d CS = −0.39 eV). It should be noted that, due to the small CSs of the components in both In peaks, the analysis was more intricate since small changes in CS and peak width/shape led to significant variations in spectral weight. The analysis of the In 3d5/2 peaks may be potentially further complicated by the observation of a peak asymmetry for clean InGaAs with surface O below the XPS detection limit.50 The asymmetry was attributed to core hole screening due to free electrons and may also thus potentially apply to InP. However, the InP band gap is much larger than that of InGaAs, and the Fermi level was close to midgap, as discussed below, indicating that the free carrier density should be rather small. Such an asymmetry can in principle be modeled using Doniach−Sunjic line shapes51 or synthetic line shapes including an exponential tail function. For the as-cleaned surfaces, both the In 3d5/2 and In 4d peaks showed a clear shoulder at higher binding energies (see Figure 3a,f). Such spectra could not be fitted by an asymmetric line alone (nor by Doniach−Sunjic line shapes or an exponential tail function) nor in combination with a second component representing In0. This clearly demonstrates the necessity of an In3+ component to correctly account for the experimental spectra. In keeping with the above analysis of the P 2p peak, the CSs of the In3+ components were inconsistent with the reported values for the InPO4 orthophosphate and metaphosphate40 phases, which are on the order of 1 eV, but were consistent with the values for Inx(HPO4)y.49 It should also be noted that the CSs were similar to values observed for In2O3;40,52 hence, an admixture of In2O3 is also consistent with the data. We now turn to the quantification of the InP native oxide thickness. Because of the small effective attenuation length of

Figure 4. Relative depth of the different components (a) of P, defined as the ratio of the spectral weights of P 2p at hν = 250 and 660 eV, and (b) of In, defined as the ratio of the In 3d5/2 and In 4d spectral weights. A higher value indicates that the component is closer to the surface.

of the spectral weights of the components at hν = 250 eV to those at hν = 660 eV. A larger value thus signifies that the component is closer to the surface. The results indicate that the outermost surface consisted of the P(2±Δ)+ suboxide with the P5+ component underneath. By contrast, P0 was found beneath the phosphorus oxides at the interface with InP. These findings are consistent with previous reports that the surface oxidation of InP proceeds by 1081



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(not shown) were consistent with an unreconstructed (1 × 1) InP covered by an amorphous oxide. The observation of exclusively monolayer steps suggests a common termination of the InP beneath the surface oxide on all terraces. It has been reported that the In termination of InP is much more stable in aqueous solutions because of the much higher solubility of phosphorus oxides.58 This is also consistent with the observation of a slightly In-rich surface oxide, as discussed above. Exposure to TMA. Both the P 2p and In 3d5/2 photoelectron spectra showed significant changes upon exposure to TMA (Figures 1 and 3). The variation of the spectral weight of the different components in P 2p (at both 250 and 660 eV), In 3d5/2, and In 4d is summarized in Figures 1d,h and 3d,h, respectively. The most prominent change in the P 2p spectra was the appearance of a novel component at higher BE. The CS of this component increased very slightly from 5.91 to 6.03 eV (with respect to the P−In substrate component) from one to four TMA pulses. This was accompanied by the (partial) disappearance of the P5+ component (CS = 4.20 eV). This suggests that TMA reacts with the surface P5+ by forming a compound which is not present on the initial surface after cleaning and degassing. In addition, we also observed a reduction of the P(2±Δ)+ suboxide. By contrast, the intensity of P0 was only very slightly affected, and P0 was found to persist after TMA exposure. Previously, the novel component in P 2p has been attributed to P2O5.33 It should however be remarked that the CS of P2O5 was reported to be 6.8−7.0 eV with respect to the peak for the InP substrate,40,41 significantly larger than the observed values (both in this work and in ref 33). In addition, the formation of P2O5 is unlikely because of several thermodynamic instabilities. First, P2O5 is not thermodynamically stable on InP.25,26 As a result, P2O5 has been observed to react with InP at temperatures comparable to the ALD process temperatures,59 forming InPO4 (or In(PO3)3). Furthermore, the formation of P2O5 by TMA exposure would have to be concurrent with the formation of Al2O3. If we assume for simplicity that the surface oxide consisted approximately of InPO4, P2O5 could in principle be formed according to

In and P outdiffusion rather than O indiffusion. The diffusion of P is considerably slower than that of In, leading to a P0 enrichment at the interface between InP and the surface oxide.54 Relative depth information of the In components can be obtained from a comparison of the spectral weights in In3d5/2 (more surface sensitive) with respect to In 4d (less surface sensitive) at the same photon energy of hν = 660 eV. The results indicate that In0 was closest to the surface. This may indicate the presence of small metallic In clusters at the surface. As expected, In3+ was also closer to the surface than the substrate InP. On the basis of the relative depth plots, the calculated effective attenuation lengths,53 and the tabulated cross sections,55 we can estimate the amount of the surface oxide. A highly accurate determination of the composition profile of the surface region of the sample is beyond the scope of this paper. However, a semiquantitative assessment was possible using a simplified model based on the relative depth plots. For simplicity, we limit ourselves to a four-layer model with the InP substrate underneath, followed by a P0 layer, the P5+ “surface oxide” layer comprising also the In3+ component, and finally a surface layer comprising the P (2±Δ)+ , S1, S2, and In 0 components.56 Within this model, the amount of P0 present at the interface was on the order of 0.3 Å, whereas the combined surface layer was approximately 1 Å. The thickness of the surface P5+ (hydrogen) phosphate layer was about 4.5 Å and appeared slightly In-rich, which may indicate the additional presence of In2O3. The Fermi level position within the InP band gap could be deduced from the valence band edge spectrum in Figure 2b. The Fermi level was found approximately 600 meV above the valence band edge. This is much higher than what would be expected for p-type InP under flat band conditions and indicates Fermi level pinning close to the midgap position. This falls within the previously reported range for p-type InP5,57 and has been attributed to P vacancies and/or P antisites.5 Figure 5a shows a scanning tunneling micrograph of the InP surface after ex situ cleaning in H2SO4 and rapid introduction

2InPO4 + 2Al(CH3)3 → Al 2O3 + P2O5 + 2In(CH3)3 (1)

However, the concurrent formation of P2O5 and Al2O3 is unlikely considering the thermodynamics of the system. This is illustrated by the calculated Al2O3−P2O5 pseudobinary diagram at typical ALD process pressures60 in Figure 6. The phase diagram indicates that AlPO4 (or mixtures of AlPO4 with either Al2O3 or P2O5, depending on the precise surface stoichiometry) rather than a mixture of Al2O3 and P2O5 should be formed at the relevant process temperatures on the surface. As a matter of fact, formation of AlPO4 from Al2O3 and P2O5 is strongly exothermic with a Gibbs free energy gain of ∼290 kJ/mol. Furthermore, at ALD process pressures and temperatures, P2O5 is expected to be gaseous (see also Figure 6). Hence, the observation of P2O5 on the surface after TMA exposure is unlikely, and the novel component in the P 2p spectra should be attributed to AlPO4, in accordance with the bulk phase diagram. In contrast with eq 1, AlPO4 could be formed schematically according to

Figure 5. Scanning tunneling micrographs of the InP surface (a) after H2SO4 cleaning and (b) after subsequent exposure to TMA. The inset in (b) shows an expanded view of the surface.

into a vacuum. The surface was characterized by atomically flat terraces separated by monolayer-high steps (step height a/2 ≈ 2.9 Å) along the [110] direction. The rms roughness was 2.1 Å. No particular structure could be resolved on the terraces, which suggests the presence of a disordered surface due to oxidation. In keeping, low-energy electron diffraction (LEED) patterns 1082



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In the In 3d5/2 and 4d spectra, exposure to TMA led to a reduction of the spectral weight of the In3+ component in combination with a shift to lower binding energies. By contrast, a small relative increase in In0 is observed, which may be due to In0 formation during the surface reaction with TMA or to In0 segregation. As already mentioned above, this analysis was complicated by the small CSs of In3+ with respect to In−P. An unambiguous quantitative separation between intensity reduction and CS variation was thus difficult since both were strongly correlated, yet satisfying fits of both In 3d5/2 and In 4d within a common scenario with consistent peak widths and shapes could not be obtained when only the In3+ intensity or the BE was allowed to vary. This indicates that both effects did occur during TMA exposure. The absolute quantification of the evolution of the In3+ spectral weight was however hampered by the correlations with the CS variation; the results should thus be considered as semiquantitative only. Results for best fits of In 3d5/2 and In 4d are shown in Figure 3d,h, respectively. It should be mentioned that the inclusion of Doniach−Sunjic or synthetic line shapes including an exponential tail function could not account for the spectra after TMA exposure without the presence of an In3+ component. While this led to a quantitative difference in the spectral weights, the qualitative scenario was robust. The results thus indicate that the deoxidation was only partial, with most of the In3+ spectral weight remaining after TMA exposure. We speculate that the shift of the In3+ BE may be related to the phosphorus (and/or hydrogen) transfer from In to Al, leading to more In2O3-like behavior, which should lead to a smaller CS (cf. also, e.g., the CSs of oxides vs hydroxides of Al48). The relative depth plots of both P and In components after TMA exposure are also shown in Figure 4. As expected, the AlPO4 component was found at the very surface. Generally, very little change was observed for all other components. A notable exception was the elemental P0 component, which was found at the interface with InP after H2SO4 cleaning but appeared to segregate to the surface after TMA exposure. This is consistent with the P+ ToF-SIMS profiles in Figure 7, which

Figure 6. Calculated Al2O3−P2O5 pseudobinary phase diagram at a pressure of 5 mbar. (s) denotes solid phases, whereas (g) denotes gaseous phases. The dotted line indicates the temperature used in our experiments.

InPO4 + Al(CH3)3 → AlPO4 + In(CH3)3

(2)

The P 2p CS of bulk AlPO4 with respect to InP has been reported to be 5.8 eV,40,61 which is in excellent agreement with the experimental observations of 5.9−6.0 eV. Hence, from the experimentally observed CS, the novel peak was fully consistent with the formation of AlPO4. The formation of a surface phosphate compound was further corroborated in the valence band spectra in Figure 2a. While such features were absent after degassing, two peaks around 11.4 and 15.4 eV appeared upon TMA exposure. Such features have previously been linked to the formation of phosphates and cannot be attributed to pure oxides.47,48 The doublet splitting and the peak position are consistent with the formation of an AlPO 4 aluminum orthophosphate.61 Note that for meta- or polyphosphates or P2O5, the doublet should show different splittings (∼5 eV for the Al4(P4O12)3 polyphosphate and ∼2.5 eV for P2O5) as well as slightly different peak positions.62 Reaction 2 corresponds to a “ligand exchange” mechanism, similar to what was previously suggested for the reduction of native group III oxides on III−V arsenide surfaces by TMA.14,17 However, the ligand exchange occurs with the phosphate and does not reduce P5+, as observed for high-valence arsenic oxides during the TMA exposure of GaAs.17 In principle, P5+ (as in PO43−) could react with methyl groups by forming trimethyl phosphate or, after the reduction of P5+, trimethyl phosphite or trimethylphosphine. All these compounds are volatile at the process temperature and should thus lead to a removal (rather than a transformation) of the high-valence phosphorus oxides. Hence, the data indicate that these reactions are not favored, presumably because of the high stability of AlPO4. The above discussions neglect the potential presence of additional H, as in Inx(HPO4)y. The presence of H may lead to the formation of CH4 concurrently with metallic In0. Such a process is also consistent with the increase in In0 during the exposure to TMA (see below), although the increase was only minor, indicating that this may not be the principal reaction path. It is also possible that an Alx(HPO4)y hydrogen phosphate was formed. This is difficult to establish from SRPES because of the lack of standards.48 It should however be remarked that such reactions may only describe very schematically the real surface reactions and further detailed theoretical studies are required to fully understand the surface chemistry during TMA exposure of InP (see, e.g., ref 63 for potential reaction pathways of the TMA/gallium and arsenic oxide chemistry).

Figure 7. P+ ToF-SIMS profiles of Al2O3/InP. Prior to Al2O3 ALD, the InP surface was cleaned ex situ in H2SO4 or (NH4)2S as indicated. More P was found in Al2O3 after H2SO4 cleaning due to increased P segregation.

showed that more P extended further into Al2O3 after H2SO4 cleaning than after (NH4)2S passivation. As shown in the next section, the (NH4)2S passivation reduced the amount of P0 on the initial surface below the SRPES detection limit and can thus be used as a reference. The Al 2p SRPES spectra after TMA exposure are depicted in Figure 8. No Al was detected on the starting surface. The 1083



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after ex situ treatment in (NH4)2S and degassing at 300 °C are depicted in Figures 9a and 10a, respectively. The P 2p spectrum

Figure 8. (a) Al 2p SRPES spectra (photon energy hν = 640 eV) after H2SO4 cleaning and degassing at 300 °C as well as after exposure to TMA as indicated. The data are offset for clarity. (b) Evolution of the total Al 2p yield of the components as a function of the number of TMA pulses. In all cases, the Al 2p photoelectron yield was normalized to the total electron yield.

Figure 9. P 2p SRPES spectra (photon energy hν = 250 eV) (a) after (NH4)2S cleaning and degassing at 300 °C as well as (b) after exposure to TMA. The spectra were fitted with the five components described in the text. (c) shows the evolution of the relative spectral weight of the components as a function of the number of TMA pulses.

total Al 2p photoelectron yield in Figure 8b shows that the TMA adsorption quickly saturated after the first TMA pulse. This indicates that the reaction with the InP surface saturated much faster than for GaAs,17 which is consistent with the observed much thinner surface oxide on InP. The deconvolution of the Al 2p spectra was complicated by of the lack of clear features in the peaks and the rather small CSs between prospective components.61,62 However, the fwhm of the peak (∼2.0 eV) was larger than what is typically observed on ALD-grown Al2O3 (∼1.2 eV). This suggests the presence of multiple components in the spectrum, yet an unambiguous deconvolution, e.g., into Al2O3 and AlPO4 components, was not possible because of the lack of sufficiently accurate standard measurements and the reported small BE differences61,62 of less than 1 eV. Although there was little shift of the peak maximum, the spectral weight at higher BEs increased slightly between one and four TMA pulses, rendering the peak more symmetric. This may indicate stronger oxidation of Al for prolonged TMA exposure, similar to what has been reported for GaAs.17 An STM micrograph of the InP surface after TMA exposure is shown in Figure 5b. The terraces, separated by InP monolayer steps, which were observed on the H2SO4-cleaned substrate, still remained visible after TMA exposure. In addition, small clusters with average diameters of 3.7 ± 1.0 nm and heights of 0.5 ± 0.2 nm appeared. The cluster density was 5 × 1012 cm−2, about 2 orders of magnitude below the saturation density of TMA adsorption for steady-state ALD (i.e., on OH-terminated Al2O3 surfaces). These clusters can thus rather be related to the formation of AlPO4. This shows that the reacted AlPO4 did not form a uniform continuous layer but rather islandlike clusters with a lateral dimension of a few nanometers, possibly because of nucleation and diffusion kinetics. Similar to GaAs,17 the Fermi level was found to move only very weakly upon TMA exposure. A shift of ∼70 meV toward the conduction band edge was observed after the first TMA pulse (Figure 2b) with no effect of further TMA exposure. This suggests that the interfacial defect density between InP and AlPO4 was not reduced below levels in the 1012 cm−2 range.23 The shift toward a pinning position even further from the valence band edge may stem from an increase in interfacial defect density or from a change in the defect energy position within the band gap upon TMA exposure. TMA Exposure of (NH4)2S-Passivated InP. InP Surface after (NH4)2S Passivation and Degassing. In addition to the H2SO4 cleaning of the InP surface, the wet (NH4)2S passivation of InP was also examined. The P 2p and In 3d5/2 spectra of InP

Figure 10. (a, b) In 3d5/2 and (c, d) In 4d SRPES spectra (photon energy hν = 660 eV). Spectra both after (NH4)2S cleaning and degassing at 300 °C (a, c) and after exposure to TMA (b, d) are shown. The spectra were fitted with the three components described in the text.

could be accounted for by the same set of components as for the H2SO4-cleaned InP. The major difference between the two wet surface treatments was the reduced spectral weight of the high-valence P5+ component. The CS of this component after (NH4)2S passivation (4.59 eV) was slightly larger than for the H2SO4 cleaning (4.20 eV), possibly because of the influence of (second neighbor) S atoms at the interface or a reduction of the hydrogen phosphate formation. In contrast with the H2SO4-cleaned InP surface, the P(2±Δ)+ suboxide and P0 components were below the detection limits after (NH4)2S passivation. Also, no PxSy mono- or polysulfide bonds could be detected, since they would lead to peaks with BEs between 130 and 132 eV.64,65 As a result, the (NH4)2S treatment clearly reduced the surface oxidation both in solution and/or in air, although the protection did not appear complete. The In 3d5/2 and 4d spectra in Figure 10 indicated the presence of an In3+ component on the surface, as well as metallic In0, similar to the spectra after H2SO4 cleaning. The In3+ CSs with respect to the peaks for the InP substrate component (0.38 eV for 3d5/2 and 0.42 eV for 4d) were almost 1084



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identical to the values obtained for H2SO4-cleaned InP. It should be noted that the CSs of In2O3 and InS (and Inx(HPO4)y) are essentially indistinguishable,66,67 which renders it difficult to separate surface oxides, hydrogen phosphates, and sulfides. However, as for H2SO4-cleaned InP, phosphate phases such as InPO4 or In(PO3)3 with CSs on the order of 1 eV were not observed. This was again consistent with the measured valence band structures around 15 eV, as shown in Figure 11a and discussed above. Figure 12. S 2p SRPES spectra (photon energy hν = 690 eV) (a) after (NH4)2S cleaning and degassing at 300 °C as well as (b) after exposure to TMA. The spectra were fitted with the three components described in the text.

layers below the surface, and (3) polysulfidic S on the surface, either as surface dimers bonded to two In atoms or possibly in the form of small S clusters (component C, BE = 162.4 eV). In all three cases, S was bonded to In (as well as to S for component C) and not to P, in agreement with the absence of P−S bonds in the P 2p spectra. The data indicate that S was present in all three binding configurations after the wet (NH4)2S passivation. If one takes the photoelectron attenuation into account, the majority of S was bound in subsurface sites with an additional sizable amount in bridge positions. By contrast, only ∼15% of the spectral weight corresponded to the surface dimers (or clusters). After the (NH4)2S passivation, the Fermi level position was found about 850 meV above the valence band edge (Figure 11b) and thus shifted even further toward the conduction band edge with respect to that of H2SO4-cleaned InP. This indicates that the p-InP surface was weakly inverted. The Fermi level position in the upper half of the band gap may be understood by surface electron doping by S, in particular in subsurface or bridge sites. Indeed, S is known as a bulk n-type dopant in InP when substituting P,74 yet the Fermi level position was still well below the conduction band, suggesting that a rather large number of surface acceptor states exist in the midgap region. Thus, it appears that the reduction of the surface oxide by (NH4)2S passivation was not necessarily associated with a strong reduction of the surface defect density. The STM micrograph in Figure 13 shows that the surface was corrugated on small length scales with terrace/island dimensions of only a few nanometers at most, unlike the large terraces obtained after H2SO4 cleaning. Nonetheless, the typical corrugation height was only around 1−2 monolayers (5−10 Å),

Figure 11. (a) Valence bands and (b) a close-up of the valence band edges of (NH4)2S-passivated InP after degassing as well as after exposure to TMA. In (a), the data are offset for clarity. The photon energy was hν = 150 eV.

For both In 3d5/2 and 4d, the relative spectral weights of the three components indicated a slightly weaker In3+ component on (NH4)2S-passivated InP (spectral weight ∼0.35) with respect to H2SO4-cleaned InP (spectral weight ∼0.40−0.45). However, this reduction was much smaller than for the P5+ component in P 2p. This suggests that a significant part of the In3+ spectral weight may stem from In−S bonds if one assumes that the In/P ratio in the surface oxide did not vary strongly. This is also consistent with the discussion of the S 2p peak below. By contrast, the spectral weight of the In0 component on the two InP surfaces was identical within the accuracy of the deconvolution. Again, it has to be noted that the quantification of the amount of In3+ and In0 on the surface was less accurate because of the small CSs (as well as the potential peak asymmetry), as already discussed above. The layer structure of the InP surface oxide, as determined from relative depth plots (not shown) was very similar to that of H2SO4-cleaned InP. An equivalent three-layer model led to ∼2 Å of surface P5+/In3+ (for comparison, this was ∼4.5 Å after H2SO4 cleaning). The total combined thickness of the layer containing the P surface components and In0 was again ∼1 Å. The wet (NH4)2S passivation of InP also led to the presence of S on the surface, as indicated by the S 2p spectra in Figure 12. The initial spectrum after degassing could be accounted for by three components. All components had BEs in the range expected for sulfides (S2−), and no S−O sulfate components were observed in the spectra, indicating that S was not oxidized during the short air exposure.68,69 Very similar spectra and deconvolutions have been obtained previously by SRPES,70−72 although their interpretation has not been unanimous. A thorough investigation of the S sites corresponding to the different components has been performed by the azimuthal dependence of the SRPES spectra71 in combination with structural71 and ab inito73 modeling. According to these studies, the three components can be attributed to (1) S in a bridge position on the surface bonded to two In atoms (component A, BE = 161.1 eV), (2) subsurface S (component B, BE = 161.9 eV) on substitutional sites (SP), replacing P in the topmost

Figure 13. Scanning tunneling micrograph of the InP surface after (NH4)2S passivation. The full color scale corresponds to 25 Å, and the typical corrugation amplitude is 5−10 Å. 1085



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Figure 14. (a−e) P 2p SRPES spectra (photon energy hν = 250 eV) after individual ALD (sub)cycles on H2SO4-cleaned InP as indicated. The spectra were fitted with the seven components described in the text. (f) and (g) show the evolution of the relative spectral weight of the components and the AlPO4 chemical shift, respectively.

AlPO4 occurred independently of the surface preparation and also independently of details of the initial surface chemistry. Since the amount of P5+ was lower on (NH4)2S-passivated InP than on H2SO4-cleaned InP, the amount of AlPO4 was also reduced. The remaining P5+ component at a CS of 4.59 eV was reduced close to the detection limit after TMA exposure, indicating the surface reaction was much more complete than on H2SO4-cleaned InP. It should be noted that further TMA exposure did not lead to detectable changes in all spectra. This is consistent with the lower amount of surface oxide, which means that less TMA is necessary to saturate the surface reactions. By contrast, the In 3d5/2 and In 4d spectra (Figure 10b,d) showed only a very weak spectral weight reduction of the In3+ component. This may be due to the fact that a significant part of this component was due to In−S bonds, in particular subsurface bonds. These subsurface bonds were not affected by the TMA exposure, as demonstrated in the discussion of the S 2p spectra below. As in the case of H2SO4 cleaning, the spectral weight of In0 increased very weakly, either due to In0 formation or as the result of surface segregation. Figure 12b shows the effect of the TMA exposure on the S 2p spectra. Again, only sulfide-like components were visible with no traces of oxidized sulfite or sulfate. Not unexpectedly,

and the resulting rms roughness was 2.8 Å, rather similar to the value obtained for H2SO4-cleaned InP (2.1 Å). It should be noted that both the magnitude and the lateral length scale of the roughness were very similar to what was observed on asreceived InP wafers (data not shown). This suggests that the (NH4)2S passivation does not necessarily deteriorate the surface roughness much, at least for the conditions used in this study, somewhat in contrast with what has been observed, e.g., for III−V arsenides.75 LEED (data not shown) exhibited an unreconstructed (1 × 1) InP pattern, indicating the absence of clear S-induced surface reconstruction. This is in contrast with previous reports79 but may be due to a different annealing temperature since it was found that S-induced (2 × 1) reconstructions were only formed above typical ALD process temperatures. We remark that the lack of a surface reconstruction is not necessarily linked to surface disorder, as indicated by the small-scale surface roughness since similar surface morphologies were found by STM on reconstructed surfaces also.76 Exposure to TMA. When the (NH4)2S-passivated surface was exposed to TMA, the surface reactions were similar to what was observed on H2SO4-cleaned InP. In the P 2p spectrum (Figure 9b), the P5+ component experienced the same shift as for H2SO4-cleaned InP. This indicates that the formation of 1086



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the subsurface SP component B was not much influenced by TMA exposure. By contrast, the surface bridge site component A was much reduced, indicating that most of the surface S was removed during the TMA exposure. The surface dimer component C was relatively unaffected, indicating a higher dimer stability, which is consistent with ab initio calculations.73 This indicates that some S indeed remained present at the interface between AlPO4 and InP after TMA exposure, although the majority of the residual S occupied subsurface sites, bonded to In. Hence, while S may passivate the surface of InP, it does not necessarily passivate the interface between InP and AlPO4/Al2O3. The presence of S in the stack did not have any marked influence on either CSs or the absolute BEs of the P 2p or In 3d5/2 spectra. This indicates that S did not lead to a measurable dipole in the structure. This is in good agreement with measurements by internal photoemission, which showed no marked influence of (NH4)2S on the band offsets77 and were fully consistent with little residual S present at the interface. The removal of a large part of the surface S from the structure led to a movement of the Fermi energy toward the valence band edge, as shown in Figure 11b. This is consistent with a reduction of the surface electron doping by S. The Fermi level position after TMA exposure (∼650 meV above the valence band edge) was very similar to that obtained after TMA exposure of H2SO4-cleaned InP. While a direct correlation between Fermi level position and interface defect density is difficult, this result nevertheless suggests that the (NH4)2S passivation of the surface did not significantly reduce the interface defect density in final Al2O3/AlPO4/InP stacks. Further Atomic Layer Deposition of Alumina. We will now discuss the effect of subsequent H2O pulsing and further TMA/H2O cycles on the chemistry of the interface with InP. The P 2p and In 3d5/2/In 4d spectra after subsequent H2O and TMA exposures beyond the first TMA pulse(s) are shown in Figures 14 and 15, respectively. All spectra were obtained on H2SO4-cleaned InP; no qualitative differences in the P 2p and In 3d5/2/In 4d behavior were found for the ALD overgrowth for (NH4)2S-passivated InP (data not shown), so the discussion below also applies to (NH4)2S-passivated InP, with the only differences being the thickness of the AlPO4 interfacial layer and the absence of detectable P0 segregation. At first sight, the P 2p spectra remained rather unchanged, which suggests that further ALD essentially deposited Al2O3 on the surface and did not significantly alter the interface formed during the first TMA exposure. In particular, no significant additional AlPO4 formation or any further oxidation of the InP or the interface layers was observed after H2O exposure and subsequent ALD cycles (see Figure 14f). The relative depth plots of the different components in P 2p were also consistent with the overgrowth of the interfacial AlPO4. P0 remained close to the surface, indicative of continuing P segregation. This is in agreement with the ToFSIMS profile in Figure 7, although it is difficult to deduce the relative position of P0 with respect to the deposited Al2O3 from SRPES. An interesting observation was that the CS of the AlPO4 component with respect to the InP substrate oscillated between higher values after TMA exposure and lower values after H2O (Figure 14g). Such an oscillation can be attributed to the formation of surface OH dipoles after reaction with H2O. The surface dipoles lead to the formation of an image dipole at the interface between the Al2O3/AlPO4 dielectric and the semi-

Figure 15. (a−e) In 3d5/2 SRPES spectra (photon energy hν = 660 eV) after individual ALD (sub)cycles on H2SO4-cleaned InP as indicated. All spectra were fitted with the three components described in the text. (f) Evolution of the spectral weight of the three components in In 3d5/2.

conducting InP. Hence, only spectral components originating from the Al2O3/AlPO4 layer experience a BE shift, whereas the BE of substrate components will remain fixed. This then leads to the observed CS variations between P5+ and P−In. Indeed, both the sign and the magnitude of the CS oscillations were consistent with typical values for the OH dipole moments (∼1 D), OH surface densities,36 and the dielectric constant of the Al2O3/AlPO4 layer. Similar to the P 2p spectra, the In 3d5/2 (Figure 15) and In 4d (not shown) SRPES spectra showed only negligible changes upon further ALD. The same was observed for relative depth plots (not shown). This is again consistent with the Al2O3 overgrowth of the interfacial layer formed by TMA exposure. No components with CS oscillations were found in the spectra, which may suggest that the remaining In3+ component was not located within the Al2O3/AlPO4 layer but at the interface with InP. This is consistent with the discussion of the TEM results below. However, the small CSs and the strong correlations between the CS and spectral weight of the In3+ render a definitive conclusion rather difficult. The Al 2p spectra in Figure 16a showed an increasing intensity as a function of the number of ALD cycles, as expected. An interesting feature was again the oscillatory behavior of the apparent peak maximum, which was found to shift to higher BE after TMA exposure and to lower BE after H2O exposure. Note that the sign of the CS is incompatible with being caused by a varying contribution of Al−CH3 bonds to the Al 2p spectra. Although no binding energies for Al−CH3 bonds on Al2O3 surfaces have been reported, electronegativity 1087



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The valence band edge spectra in Figure 17 showed no evolution during further ALD, either on H2SO4-cleaned (Figure

Figure 17. Evolution of the InP valence band edge (photon energy hν = 150 eV) during the ALD of Al2O3 on (a) H2SO4-cleaned and (b) (NH4)2S-passivated InP. No variation in the Fermi level position was observed on both initial InP surfaces. Figure 16. (a) Al 2p SRPES spectra (photon energy hν = 640 eV) after individual ALD (sub)cycles on H2SO4-cleaned InP as indicated. (b) Evolution of the Al 2p peak position during ALD. (c) Al2O3 growth curve on H2SO4-cleaned InP: atomic Al area density, as determined by ICP-MS, as well as SRPES total Al 2p photoelectron yield vs the number of TMA/H2O ALD cycles. In all cases, the Al 2p photoelectron yield was normalized by the total electron yield.

17a) or on (NH4)2S-passivated (Figure 17b) InP. In both cases, the Fermi level was pinned about 600 meV above the valence band edge. This suggests that no major modification of the interface state density or spectral distribution occurred beyond the first TMA exposure. Hence, the interface defects were determined essentially during the first TMA exposure. This also demonstrates that the (NH4)2S passivation did not lead to a different pinning position after deposition of Al2O3 by ALD as compared to H2SO4 cleaning, suggesting that the types of dominating defects were similar. It should still be mentioned that this does not necessarily rule out differences in the interface defect distribution away from the pinning position, which may strongly depend on the surface preparation and still strongly influence MOS device properties. Figure 18 shows the evolution of the S 2p peak during further ALD of Al2O3 on (NH4)2S-passivated InP. S was still found in the structure after three ALD cycles without any signs of S oxidation during the H2O pulses. A quantitative assessment of further S removal during ALD is difficult because of the small photoelectron attenuation lengths, which will render the intensities very sensitive to the position of the S components in the full stack. A close analysis of the spectra indicated an oscillatory behavior of the spectral weight of the different components, depending on whether the deposition was halted after a TMA or H2O pulse. After H2O exposure, the surface dimer component (polysulfidic S, component C) clearly increased with respect to the surface/interface bridge S (component A) and the subsurface S (component B). In addition, there was an increase in the spectral weight of component C with the number of ALD cycles. We remark that the behavior could not be well accounted for by keeping the spectral weights constant and allowing the binding energies (especially that of component C) to vary. Hence, the behavior was distinct from the CS oscillations observed for P5+ in P 2p and the peak maximum in Al 2p. By contrast, the overall evolution may be explained by the surface segregation of the polysulfidic S during H2O exposure. Although the underlying chemistry is not yet clear, this explains both the oscillatory behavior and the slow overall increase of the spectral weight of component C during the ALD of Al2O3. Finally, Figure 19 shows cross-sectional TEM images of a thick (10 nm) Al2O3 film on H2SO4-cleaned InP. In both the high-resolution (Figure 19a) and the HAADF (Figure 19b) images, the interface appears abrupt and regular with atomic steps clearly visible. No interfacial layer is visible. This is in

arguments suggest that the Al−CH3 groups should lead to photoelectrons at lower BEs than Al−O or Al−OH, in contradiction with the experimentally observed shifts. By contrast, both the sign and the magnitude of these oscillations are identical within experimental accuracy with the CS shift oscillations observed for P5+. This strongly suggests a common origin and can thus also be attributed to the creation and removal of surface OH dipoles by subsequent H2O and TMA pulses. Furthermore, this demonstrates that both the P5+ and (at least a considerable part of) the Al photoelectrons stemmed from the same layer in the structure and is thus a clear indication of the formation of an AlPO4 layer. Figure 16c shows the growth curve of Al2O3 on InP, as determined by ICP-MS and SRPES. Note that the reported steady-state GPC for TMA/H2O on Si substrates is 4 Al atoms nm−2,36 which was also observed for steady-state growth of Al2O3 on InP, after the film had closed and the influence of the substrate had become negligible (not shown). However, some clear deviations were visible during the first cycles. A slight growth enhancement (with respect to steady-state growth) was observed, which can be attributed to the reaction of TMA with the surface oxide. An interesting feature is the apparent growth inhibition during the second cycle, where very little Al was added to the layer. The same behavior was also found for (NH4)2S-passivated InP (data not shown). A qualitatively similar effect has been observed on Si and Ge substrates78 and attributed to the influence of the substrate on TMA absorption in the second cycle. However, the detailed mechanism of the second-cycle inhibition on InP substrates will need further consideration to determine whether the models for TMA/H2O ALD on Si and Ge substrates can also be applied for InP. During further growth, the amount of deposited Al increased with a rate similar to the steady-state GPC. The deviation between the SRPES and ICP-MS data for three cycles may stem from photoelectron absorption effects, which become non-negligible for the low photoelectron energies already at such thicknesses on the order of a few angstroms. 1088



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oxide was found to lead to an AlPO4 interlayer. This is qualitatively different from the TMA exposure of III−V arsenide surfaces, where the native high-valence oxides were removed,11−17 and can be explained by the high thermodynamic stability of phosphates (in particular with respect to arsenates). For potential MOS device applications, the presence of an AlPO4 layer is not a priori detrimental. AlPO4 is a dielectric with a large band gap of ∼7 eV (in the berlinite structure, ref 79). While the dielectric permittivity of AlPO4 is rather low (∼4.6),80 the thickness can be limited to a few angstroms and should not prohibit the scaling of the equivalent oxide thickness. The valence band offset with InP has been determined to be 2.35 eV,77 which is reasonable, although it should be noted that it is smaller than for Al2O3 (2.7 eV, ref 77). Hence, the presence of AlPO4 at the interface between InP and Al2O3 will increase the sensitivity of the MOS structure to border traps81 with respect to an abrupt interface. However, the valence band edge measurements indicate that the interface between InP and AlPO4 does not present an interface state density below the 1012 cm−2 range. Hence, the prevention of surface oxidation prior to dielectric deposition still appears to be necessary to obtain low-defect interfaces. These results show that wet (NH4)2S passivation can indeed reduce the surface oxidation of InP, however not below an equivalent thickness of about 2 Å. Because of the observation of exclusively sulfidic S, we speculate that the S does indeed effectively avoid the InP surface oxidation but the S surface density may be insufficient to prevent pinholes. In this case, the InP oxidation would be local. Future work will thus be needed to optimize the S passivation to fully suppress surface oxidation. Our results also show that the potential issue of surface electron doping by S has to be considered, which will alter MOS properties. However, we have also found that part of the S is removed during subsequent ALD, which may provide a route to limit the doping effect of the S passivation.

Figure 18. (a−e) Evolution of the S 2p SRPES spectrum (photon energy hν = 690 eV) during the ALD of Al2O3 on (NH4)2S-passivated InP. (f) shows the spectral weight of the three components, as discussed in the text.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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Figure 19. (a) High-resolution transmission electron micrograph and (b) high-angle annular dark field image of the interface between ALD Al2O3 and H2SO4-cleaned InP.

ACKNOWLEDGMENTS

We acknowledge the Helmholtz-Zentrum Berlin, Electron storage ring BESSY II, for provision of synchrotron radiation at beamline U49/2-PGM-2 and would like to thank Marcel Michling for assistance. Alexis Franquet and Laura Nyns (imec) are acknowledged for providing the ToF-SIMS data, and Olivier Richard (imec) is acknowledged for the TEM observations. We also thank Tomas Skeren (Katholieke Universiteit Leuven) for the support of the STM measurements, and Valeri Afanas’ev (Katholieke Universiteit Leuven) and Dennis Lin (imec) for valuable discussions. The research leading to these results has received funding from the European Community's Seventh Framework Programme (FP7/20072013) under grant agreement No. 226716. L.N.J.R.’s work was partially funded by a Marie Curie International Reintegration Grant (IRG) under Call No. FP7-PEOPLE-2010-RG.

agreement with the formation of an AlPO4 interface layer since the TEM contrast between AlPO4 and Al2O3 can be expected to be very small. This furthermore indicates that the In content in the AlPO4 layer was rather negligible since significant In incorporation can be expected to lead to a strong imaging contrast with respect to Al2O3, in particular in the HAADF images, although the presence of a (sub)monolayer of InxPOy or In2O3 at the interface is difficult to rule out.

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CONCLUSIONS In conclusion, we have studied the first stages of the atomic layer deposition of alumina on InP(100) using TMA and H2O. Both H2SO4-cleaned and (NH4)2S-passivated surfaces have been considered. The reaction of TMA with the native InP 1089



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ABBREVIATIONS ALD, atomic layer deposition; BE, binding energy; CS, chemical shift; GPC, growth per cycle; HAADF, high-angle annular dark field; ICP-MS, inductively coupled plasma mass spectrometry; LEED, low-energy electron diffraction; MOS, metal−oxide−semiconductor; MOSFET, metal−oxide−semiconductor field-effect transistor; SRPES, synchrotron radiation photoemission spectroscopy; STM, scanning tunneling microscopy; TEM, transmission electron microscopy; ToF-SIMS, time-of-flight secondary ion mass spectrometry; TMA, trimethyaluminum; XPS, X-ray photoelectron spectroscopy

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