ION IMPLANTATION INTO GALLIUM NITRIDE

Transcription

1 ION IMPLANTATION INTO GALLIUM NITRIDE C. RONNING, E.P. CARLSON, R.F. DAVIS Low Temperature Laboratory, Helsinki University of Technology Box 2200, FIN HUT, Finland L.D. Landau Institute for Theoretical Physics, Moscow, Russia AMSTERDAM } LONDON } NEW YORK } OXFORD } PARIS } SHANNON } TOKYO

2 Physics Reports 351 (2001) 349}385 Contents Ion implantation into gallium nitride C. Ronning *, E.P. Carlson, R.F. Davis II. Physikalisches Institut, Universita( tgo( ttingen, Bunsenstr. 7-9, D Go( ttingen, Germany Department of Materials Science and Engineering, North Carolina State University, Box 7909, Raleigh, NC 27695, USA Received November 2000; editor: D.L. Mills 1. Introduction Microstructural properties Damage accumulation Annealing procedures Redistribution Damage recovery Lattice sites Optical properties Defects Donor doping Acceptor doping Miscellaneous elements Rare earth elements Electrical properties Implantation isolation Donor doping Acceptor doping Device applications Summary, conclusions, and future work 377 Acknowledgements 379 References 379 Abstract This comprehensive review is concerned with studies regarding ion implanted gallium nitride (GaN) and focuses on the improvements made in recent years. It is divided into three sections: (i) structural properties, (ii) optical properties and (iii) electrical properties. The "rst section includes X-ray di!raction (XRD), transmission electron microscopy (TEM), secondary ion mass spectroscopy (SIMS), Rutherford Backscattering (RBS), emission channeling (EC) and perturbed γγ-angular correlation (PAC) measurements on GaN implanted with di!erent ions and doses at di!erent temperatures as a function of annealing temperature. The structural changes upon implantation and the respective recovery upon annealing will be discussed. Several standard and new annealing procedures will be presented and discussed. The second section describes mainly photoluminescence (PL) studies, however, the results will be discussed with respect to Raman and ellipsometry studies performed by other groups. We will show that the PL-signal is very sensitive to the processes * Corresponding author. Tel.: # ; fax: # address: cronnin@uni-goettingen.de (C. Ronning) /01/$ - see front matter 2001 Elsevier Science B.V. All rights reserved. PII: S (00)

3 350 C. Ronning et al. / Physics Reports 351 (2001) 349}385 occurring during implantation and annealing. The results of Hall and C}< measurements on implanted GaN are presented in Section 3. We show and discuss the di$culties in achieving electrical activation. However, optical and electrical properties are both a result of the structural changes upon implantation and annealing. Each section will be critically discussed with respect to the existing literature, and the main conclusions are drawn from the interplay of the results obtained from the di!erent techniques used/reviewed Elsevier Science B.V. All rights reserved. PACS: Vv; Jh; Ln; Ry Keywords: Ion implantation; Gallium nitride; Structural properties; Optical properties; Dopants

4 C. Ronning et al. / Physics Reports 351 (2001) 349} Introduction In the past decade, signi"cant e!ort has been applied to the emerging wide band gap semiconductors; including diamond, silicon carbide, the III}V nitrides, and their alloys. These materials are being investigated to extend the limits of device application into regimes of higher power, higher frequency, higher temperature and optical wavelength that can be achieved via the use of the mature semiconductor materials systems namely silicon and gallium arsenide. The success of the III-nitrides among the wide band gap semiconductors is due to the continuous solid solubility between gallium nitride (GaN), aluminium nitride (AlN), and the limited solubility with indium nitride (InN) and the consequent opportunity to engineer the direct band gap between 2 (red) and 6 ev (ultraviolet). Blue and ultraviolet light-emitting diodes (LEDs) as well as laser diodes have been realized [1,2] and commercialized [3] by growing structures containing multiple, high-quality, Ga Al In N layers with sharp interfaces and di!erent stoichiometries (x, y, z"0,2, 1) by metal organic chemical vapor deposition (MOCVD) onto suitable substrates. Electrons are injected during operation over a highly doped back or front layer into the active layers with lower band gap, where the electrons are "nally trapped. Spontaneous or stimulated photon emission can occur by recombination with holes. The demonstration of continuous wave operation of GaN based-laser diodes at room temperature having projected lifetimes of 10,000 h has been reported [4]. In addition, GaN-based structures o!er the prospect of superior microelectronic device performance, even in comparison to those based on other wide band gap materials. Advantages include a direct band gap, strong chemical bonding, high thermal conductivity and the realization of superior electrical properties [5]. The afore-mentioned optoelectronic devices have been achieved via MOCVD, because such a fabrication procedure is limited to relative large areas and vertical, layered device structures [6,7]. Therefore, the processing techniques used to produce new devices have to be investigated and understood, as the III}V nitride devices become more complicated [8]. Advances in microelectronics would require the possibility of adequate lateral structuring and doping techniques. Di!usion of dopants into selected areas for lateral p- and n-type doping of GaN can be excluded due to the fact that surface decomposition of GaN starts at temperatures of 8003C [9]. No signi"cant di!usion of any impurity has been observed in GaN to this temperature [10]. However, a processing technique that is widely used in semiconductor industry for lateral doping is ion implantation. Ion implantation is a convenient method to incorporate electrically and optically active dopants into the host crystal using a highly energetic beam of ions that strike and penetrate into the crystal. This method allows the introduction of a precisely controlled amount of impurity into the crystal independent of the solubility of the impurity. Ion implantation doping and isolation has played a critical role in the realization of high-performance electronic and photonic devices in all mature semiconductor material systems [11]. This is also expected to be the case for GaN and its alloys as the epitaxial material quality improves and more advanced device structures are requested. However, the main disadvantage of ion implantation is that this process is compromised by the introduction of radiation damage, that may control the optical and electrical properties and which has to be removed via annealing procedures. Even complete amorphization of the host material may occur for high implantation doses [12,13].

5 352 C. Ronning et al. / Physics Reports 351 (2001) 349}385 The following review discusses the recent advances in the success of ion implantation into GaN obtained by di!erent groups. The microstructural changes introduced by the implantation process and its annealing behavior under di!erent annealing techniques will be treated in Section 1. The resulting optical and electrical properties of ion implanted GaN with di!erent impurities obtained by a variety of characterization tools will be subsequently presented and discussed. These results will be compared with the properties of GaN doped during "lm growth. A discussion regarding necessary future research will be presented at the end of the review. 2. Microstructural properties The microstructural changes in the host substrate which occur during ion implantation are due to the loss of kinetic energy by electronic and nuclear interactions of the impinging ions with the atoms on the host lattice. In these processes, su$cient energy can be transferred to the host atoms resulting in displacements from their lattice sites. Such atoms may displace other atoms thus creating a cascade of atomic collisions. The damage caused by the impinging ions can be in the form of vacancies, interstitials, anti-site defects and extended defects, e.g. dislocations or stacking faults. The concentrations of such defects are mainly dependent on the bonding type and structure of the host material, on the mass of the impinging ion and on the implantation temperature. They can easily reach 100}1000 for each implanted impurity atom. The precise control of the impurity and damage distribution can be calculated using Monte Carlo simulations [14]; however, the exact amount and types of defects cannot be determined at this time Damage accumulation Tan et al. [15,16] demonstrated by ion beam channeling spectroscopy (RBS/C) that GaN has a high threshold for damage accumulation. A high dose of 8 10 cm Si ions implanted with an energy of 90 kev at 77 K had a negligible e!ect on the ion channeling spectrum in that it was indistinguishable from an unimplanted sample. A dose of cm Si ions was required to reach the random disorder level [15,16]. Other groups [13,17,18] have reported similar high levels for the dose threshold of damage accumulation in Si implanted GaN. Cross-sectional transmission electron microscopy (XTEM) has been used to con"rm the amorphous nature of GaN after high-dose implantation [15,16]. Damage accumulation in GaN has also been investigated using other implanted species including Ca [18}21], Ar [19,20], Mg [13,22,23] and Er [24]. Table 1 summarizes all measured critical doses for the amorphization for various species implanted in GaN [12,15,16,19,22,24}30]. The heavier ions have a lower critical dose, which is to be expected. The damage accumulation is lower for implanted Ca and Ar compared to implanted Si; however, the three ions have approximately the same mass. This e!ect may be due to dynamic annealing driven by the di!erent chemistries of the implanted impurities in GaN or simply due to the di!erent structural qualities of the GaN material used in the studies. Fig. 1 shows the relative lattice disorder (measured by RBS) as a function of the displacements per atom (DPA), that was calculated using TRIM [14], for various ions implanted into GaN. Full amorphization corresponds to 100 on the vertical scale. There is good agreement in the relative disorder caused by the di!erent implanted

6 C. Ronning et al. / Physics Reports 351 (2001) 349} Table 1 Critical dose for the amorphization of GaN for various implanted species Species Critical dose in GaN Implantation temperature Refs. (ions/cm ) (K) H ' ,26,27 O ' C * Mg ' Si , Ar Ca Er Au ; ,29 The amorphization dose has not been reached for these species at these doses and temperatures. ions compared to the calculated DPA for each ion. As seen in Fig. 1, GaN has a high threshold for damage accumulation with no signi"cant disorder occurring until a DPA of &1 is reached. Therefore, GaN has a approximately two-to-three orders of magnitude higher amorphous threshold compared to GaAs (&4 10 cm ) [31]. Conversely, the threshold for amorphization is similar to that of AlAs, where substantial dynamic annealing during implantation occurs. The addition of Al to the GaN matrix has been reported to increase its damage threshold, which is similar to the results reported for the AlGaAs system [32]. Mensching et al. [19] reported that computer simulations of the implantation induced damage yielded a factor of two higher concentration of displaced atoms than observed in RBS/C measurements. This is similar to the results of Tan et al. [15,16], who observed that GaN has a high threshold for damage accumulation and a potentially strong dynamic annealing even at liquid nitrogen temperatures. Wenzel et al. attempted to increase the dynamic annealing by performing implantation at elevated temperatures as high as 5503C [22]. However, the amount of damage increased rather than decreased with increasing temperature. A similar result was observed by Suvkhanov et al. [23], who noted that implantations performed at 7003C increased the χ minimum yield in RBS/C compared to RT or 77 K implantations. However, this is in contradiction with the results of Parikh et al. [13] and Tan et al. [33], who reported that increasing the implantation temperature decreases the damage level. The nature of the implantation damage has been investigated by several groups. We found a distinctly darker area in cross-sectional TEM images near the top surface, as shown in Fig. 2a. This darker area is attributed to contrast caused by the disorder. It is highest in the center of the damage region. High-resolution TEM images of the implanted and unimplanted areas are shown

7 354 C. Ronning et al. / Physics Reports 351 (2001) 349}385 Fig. 1. Relative lattice disorder (percentage of the aligned to random yield in the RBS-spectrum) at the damage peak as a function of displacements per atoms occurring during implantation for various ions implanted in GaN. The displacements per atom were calculated using TRIM and assuming a threshold displacement energy of 20 ev for both Ga and N sublattices. The dashed line is a sigmoidal "t to the data to help guide the eye. Fig. 2. Cross-sectional transmission electron picture of GaN implanted with 1 10 cm, 200 kev Si ions at 6503C showing the implantation region in high magni"cation (a) and high resolution (b). The high resolution image (c) shows an unimplanted region for comparison. in Fig. 2b and c, respectively. The implanted region is still crystalline with a well-de"ned atomic ordering due to the moderate implantation dose of 1 10 cm of Si ions used, but the arrows indicate small potentially amorphous pockets, which contain a high density of defects. These "ndings are in agreement with other studies, where low-dose Si-implanted GaN showed

8 C. Ronning et al. / Physics Reports 351 (2001) 349} Fig. 3. Normalized GaN (0002) XRD spectra after 180 kev Ca implantation with di!erent doses at 77 K. (From Refs. [20,21]) a near-surface region with a high density of defects including clusters, loops, and planar defects [15,16]. Similar XTEM results were obtained for low-dose Zn [34], Au [29], and C [29] implanted GaN. Examination of the energy dependence of dechanneled He ions indicate that the implantationinduced defects are mostly point defects in nature [22]. Both Hf and In implanted into GaN have heavily disturbed surroundings in the next or further nearest neighborhood after implantation, as observed by perturbed-γγ-angular-correlation spectroscopy (PAC) [35,36]. Liu et al. [20,37] found that a new peak appears at smaller angles in X-ray di!raction (XRD) spectra. Fig. 3 shows such peaks for Ca implanted GaN at 77 K (from Ref. [20,21]). The main (0002) peak of the virgin GaN-material decreases, and the new peak shifts to lower angles with increasing implantation dose. Other groups have also noticed this extra peak in GaN implanted with Mg [38,39] and Be [39,40]. This new peak is postulated to be an expansion of the GaN hexagonal lattice driven by the introduced impurities and displaced host atoms onto interstitials sites or by the incorporation of larger atoms on substitutional sites, and not a phase change in the GaN [20,41]. This explanation is supported by TEM, electron di!raction, and X-ray pole "gure measurement results [20,41]. The deconvolution of the XRD spectra reveals a better "t if three peaks are used instead of one. One peak for the virgin (0002) GaN, one peak for the expanded lattice (0002) GaN, and a third peak for an amorphous content in the GaN [20,41]. The amorphous peak arises in Ca implanted GaN at a dose of 3 10 cm, begins to dominate at a dose of 1 10 cm, then totally dominates at a dose of 6 10 cm. The onset of the

9 356 C. Ronning et al. / Physics Reports 351 (2001) 349}385 Table 2 Comparison of melting temperature and activation temperature for various compound semiconductors Melting temperature Activation temperature ¹ /¹ (3C) (3C) GaSb }600 &0.75 InP }750 &0.65 GaAs }900 &0.65 SiC } }0.57 GaN 2791?? Sublimes at ¹(¹. amorphous peak in XRD occurs at doses lower the onset of amorphous behavior, as observed using RBS/C [19,20]. This implies that the process of amorphization in GaN occurs in small local regions. These regions then increase with increasing implantation dose until the crystal collapses to create an amorphous layer. This interpretation is in agreement with TEM-results shown in Ref. [29], where planar defects seems to provide a nucleation site for amorphization. A change in the lattice parameter of GaN after implantation, similar to the extra XRD peak, has also been observed [42,43] along with a broadening of the X-ray rocking curve [34,44,45]. Furthermore, it was observed that high-dose H implantation resulted in the formation of voids or bubbles in GaN [25}27] Annealing procedures The implantation damage created in the host material can only be removed via subsequent annealing procedures. The annealing temperature for optimal implantation activation in compound semiconductors generally follows a two-thirds relationship with respect to the melting point of the material and has been intensively investigated for the common semiconductors. This behavior can be seen in Table 2, which lists the melting point of several compound semiconductors with the associated temperatures commonly reported to achieve implantation activation. The melting temperature of GaN has been determined as 27913C and [46], therefore, annealing temperatures of about 15003C are required, assuming the two-thirds relationship is applicable. Though GaN has a high melting point, it will decompose at much lower temperatures due to the very strong triple bond of molecular nitrogen (N ) that makes less negative Gibbs free energy of the nitride constituents [47,48]. The Gibbs free energy of the nitride constituents decrease with temperature faster than the Gibbs free energy of the GaN crystal resulting in decomposition before melting. Therefore, GaN surface decomposition already starts as low as 8003C [49] resulting in the formation of N and the consequent loss of nitrogen and the formation of Ga droplets on the surface [50}52]. The speed of this process depends exponentially on the temperature making it extremely di$cult to anneal GaN at high temperatures. Therefore, one has to use an annealing technique which protects the GaN surface from decomposition. In recent years, many di!erent annealing procedures have been applied or developed for implanted GaN to reach the required

10 high temperatures for annealing times in the minutes/hour range. These procedures can be divided into four groups, summaries of which are provided below. Firstly, implanted GaN samples have been annealed in nonreactive ambients. The success of this technique is limited by the product of temperature, time and pressure. Vacuum annealing o!ers the cleanest conditions; however, only temperatures around 800}9003C can be reached for several minutes without the degradation of the sample [49,53]. A slightly higher temperature can be reached under Ar or N #ow due to the higher vapor pressure, but high-purity gases are necessary to avoid the formation of thin GaO-layers on the surface. A further increase to high N overpressures (16 kbar) resulted up to annealing temperatures of 15503C [43,54]. On the other hand, the reduction of the annealing time into the second range in rapid thermal annealing (RTA) systems resulted also into maximum annealing temperatures of 12003C under N #ow [55]. Note: the measurement of the real surface temperature in RTA systems is very di$cult especially in the case of transparent samples and gas #ow conditions; thus, inaccurate values may be published. Secondly, implanted GaN has been annealed under reactive ambients [40,56}58]. In this case an exchange of nitrogen between the surface and the vapor phase occurs, i.e. the loss of nitrogen is compensated by the formation of new Ga}N bonds during the annealing process. Similar conditions exist during the growth of GaN. Therefore, these annealing techniques have been applied in MOCVD or MBE growth systems, where the implanted GaN samples are annealed under NH, an atomic nitrogen #ux or in a nitrogen plasma. Temperatures to 11003C can be reached in this way for 1 h [40,56,58]. The third way to protect the GaN surface is to deposit a capping layer that must be removed after the annealing procedure. Sputtered polycrystalline SiN and AlN have been used; however, such "lms were only stable to temperatures of 11003C [59]. The highest temperature of &13003C was achieved with epitaxial AlN-caps and subsequent vacuum annealing for 15 min [60,61]. Finally, laser processing is an additional method of annealing that uses short processing times to take advantage of GaN decomposition kinetics to suppress surface degradation. The de"nition and measurement of the equivalent annealing temperature is impossible due to the very fast nonequilibrium process. Therefore, the introduced power density is used for comparison; maximum values of 200}350 mj/cm have been applied to GaN without an observable decomposition of the surface [51,52,62]. However, this annealing technique has not yet been applied to implanted GaN Redistribution C. Ronning et al. / Physics Reports 351 (2001) 349} Most species implanted in GaN have the anticipated Gaussian distribution around the maximum concentration with a tail extending further into the GaN substrate most likely due to channeling e!ects [63}76]. The range and distribution of the implanted impurities in GaN is in good agreement with TRIM calculations [18,19,40,66]. All implanted impurities show a high thermal stability in GaN upon annealing and di!usion or a signi"cant redistribution was only observed in a few cases, as described below. Implanted hydrogen is stable in n-gan to 700}8003C; signi"cant loss is initiated at the surface at 9003C. However, the implanted H is still present after annealing at 12003C [25,26,63,64,66,67,74]. The shape of the implanted H depth pro"le is preserved constant even during the H loss, which implies that H decorates implantation-induced defects [63,64,77]. Further evidence that H is bonded to defects in implanted GaN is the fact that H can be readily di!used into as-grown GaN at

11 358 C. Ronning et al. / Physics Reports 351 (2001) 349}385 much lower temperatures [63,64,67,77]. Hydrogen implanted in p-type GaN showed a signi"cantly di!erent behavior with out-di!usion at already 5003C and a complete loss at 10003C [74]. The faster di!usion of H in p-gan is due to its signi"cantly lower formation energy with dopants, as compared to n-type material, which has been calculated by Neugebauer and Van de Walle [78}81]. Other lighter elements also show a high thermal stability in GaN, e.g., implanted #uorine does not undergo redistribution in GaN for annealing temperatures to 8003C [68] and 9003C [66]. This is surprising, as F is generally a rapidly di!using species in III}V compounds [82]. The column IV donors elements in GaN have a very high thermal stability with no redistribution for annealing to 14503C for Si [18,65,66,68,71}73,75,83}85] and 9003C for Ge [65,66,68,75,84]. Implanted oxygen, a column VI donor in GaN, showed no redistribution to 11253C [69,70,86]; however, this is in contrast to the results of Pearton et al. [87]. The last group found that O di!used into GaN from a SiO capping layer in the temperature range of 700}9003C. The rest of the column VI donors showed no redistribution during annealing at the maximum annealing temperature of 14503C [75,84,85]. Realized and potential implanted acceptors in GaN include column II and transition elements such as Be, Mg, Ca, Zn, Cd and Hg that would be metal}site acceptors and column IV elements such as C that would be N-site acceptors. Typically, acceptor species are more likely than donor species to redistribute at high temperatures in III}V semiconductors. However, similar to donor species, acceptor species in GaN appear to have a high thermal stability. Implanted Be, a common fast interstitial di!using species due to its small size, is thermally stable for anneals to 12003C [40,65,66,68,73,84,85] with slight initial broadening beginning at 9003C [73,84,85]. The remaining column II elements of Mg [65,66,68,73,76,84,85] and Ca [69,70,72] have also a high thermal stability with no redistribution observed for annealing to 14503C or 11253C, respectively. However, there have been some reports of implanted Mg redistribution at 11503C [18,71,72]. The transition element Zn has shown the highest redistribution rate of the acceptor elements implanted in GaN. Wilson et al. reported that Zn was only stable to 7003C [66] with some redistribution observed at 8003C [65,68]. Strite et al. observed Zn di!usion at 1100}11503C [34,43] with rapid di!usion at 12503C [43] such that the Zn di!used throughout the GaN "lm. These results were con"rmed by Suski et al. [54]. The column IV acceptor carbon is stable up to the maximum annealing temperature of 14503C [65,68,75,84,85]. Several rare earth elements have been implanted in GaN including praseodymium (Pr), neodymium (Nd), europium (Eu), holmium (Ho), erbium (Er), thulium (Tm), ytterbium (Yb), lutetium (Lu), thorium (Th), and uranium (U) [68]. These elements redistribute at lower temperatures with increasing implantation dose. No redistribution was found to annealing temperatures of 8003C for doses less than 2 10 cm ; however, the onset of redistribution drops to 7003C for doses higher than 5 10 cm [68] Damage recovery The recovery of structural implantation damage can be directly observed by RBS/C, XRD, TEM and PAC. The "rst two techniques are mainly sensitive to the crystal structure; whereas, the other two are able to detect point defects. The results of all these measurement techniques show no recovery of the accumulated damage in GaN after post-implantation annealing for low

12 C. Ronning et al. / Physics Reports 351 (2001) 349} Fig. 4. XRD 2-axis ω-2θ map of the (0002) GaN peak (a) directly after Be implantation with a dose of cm and after subsequent annealing to a temperature of 9003C (b) for 10 min. (Note: contour plots have log scale, from Ref. [40]) Fig. 5. Aligned RBS spectra of Ge-implanted GaN with a dose of 5 10 cm and an energy of 300 kev. The aligned spectra were taken along the [0001] direction after implantation and after annealing at 9003C for 1 h. Also included are a random spectrum and an aligned spectrum of a virgin GaN sample. temperatures up to 6003C. The spectra are almost identical to the spectra taken from the as-implanted situation. Annealing of implanted GaN between 6003C and 11003C results in a reduction of the implantation damage. This can be clearly seen in Figs. 4}6.

13 360 C. Ronning et al. / Physics Reports 351 (2001) 349}385 Fig. 6. Perturbed-γγ-angular-correlation spectroscopy (PAC) results measured on In in GaN as a function of annealing temperature. (From Ref. [60]) Fig. 4 shows that the implantation correlated extra peak in the XRD-pattern can be completely eliminated after annealing to 9003C in GaN implanted with a low implantation dose ( cm ) of Be [38]. This indicates that the induced structural damage can be completely removed; however, a reduction and not a complete elimination of this peak is observed for heavier elements and higher implantation doses, even for higher annealing temperatures [39,40,54,56]. Fig. 5 shows RBS/C spectra taken from a high-dose Ge-implanted GaN sample, which was annealed to a temperature of 9003C. The onset of recovery is visible after the annealing step, but the near surface implantation damage is still present. The height and reduction of this damage peak strongly depends, similar to the XRD-results, on the dose and annealing temperature and is in agreement with other published studies [17}20,22,32,33,85,88}91]. A complete elimination of the damage peak was only observed for low implantation doses and very high annealing temperatures [32]. Plan view TEM con"rmed that annealing at 11003C for 30 s did not fully recover the damage [83,85,92]. For high implantation doses beyond the amorphization threshold, it was furthermore found by XTEM that the implanted region recrystallized into a polycrystalline layer with no detectable epitaxial re-growth [33,88]. The most sensitive technique to monitor the direct surrounding of the implanted impurities is PAC [93]. Fig. 6 shows the fraction of implanted In atoms with a defect-free neighborhood as well as with disturbed surroundings as a function of annealing temperature [35,60]. For isochronal annealing treatments in vacuum a gradual recovery of the implantation damage in the surrounding volume of each indium atom occurs between 6003C and 9003C. After annealing at 12003C, approximately 70(5)% of the probe atoms occupy undisturbed sites, and the remaining fraction of indium atoms occupy sites with weakly disturbed surroundings [35]. Therefore, point defects remain in the GaN sample even after this high annealing temperature. These "ndings were con"rmed by PAC measurements after ion implantation of Hf into GaN [36]. Around 65% of the Hf occupy defect free sites after annealing at 9003C Lattice sites Implanted impurity atoms can occupy several site locations in the host lattice including substitutional and interstitial lattice sites. The resulting electrical and optical properties of the

14 C. Ronning et al. / Physics Reports 351 (2001) 349} implanted GaN sample strongly depend on the local bonding con"guration and thus the lattice site location of the impurity atoms. The lattice site of the impurity may change upon post implantation annealing procedures, also a!ecting the electrical and optical properties. Therefore, the knowledge of the lattice sites of the implanted impurities is essential for device processing by ion implantation. Several groups have used assorted channeling and spectroscopy techniques to determine the lattice site location of various impurities implanted in GaN. We used the emission channeling (EC) [94] technique for the determination of the lattice site location of implanted Li [95], Na [60], In [35], Sr [60], Tm [96], and Yb [96] in GaN. The EC-technique makes use of implanted radioactive probe atoms and measures the channeling e!ects of emitted decay particles such as conversion electrons, β-particles, or α-particles [94]. Fig. 7 shows a contour plot of a two-dimensional α-ec pattern measured around the c-axis 0002 at room temperature after implantation of Li into GaN. The emission pattern has a six-fold symmetry showing three pronounced planar channeling e!ects and three pronounced planar blocking e!ects. Furthermore, in the c-axis direction the normalized emission yield is larger than unity, indicating channeling e!ects along the c-axis direction. Channeling of α-particles can only occur if the lithium emitter atoms are located in interstitial sites with respect to the c-axis atom rows. On the other hand, blocking e!ects occur if the emitter atoms are located within an atomic row or plane. The axial and the three planar channeling e!ects are then due to emitter atoms located in between the indicated equivalent (3 120), (3 210) and (01 10) planes. The observed three planar blocking e!ects are due to emitter atoms located in the equivalent (2 110), (1 010) and (1 100) planes. The Li emitter atoms are therefore unambiguously located in the center of the hexagons. Emission channeling patterns around the c-axis were also recorded for increasing substrate temperatures to 5003C [95]. The normalized yield measured in the c-axis direction as a function of temperature is plotted in Fig. 8. The onset of Li-di!usion and a lattice site change occur in the rather narrow temperature regime of 410$253C. We observed a clear indication for a lattice site change of a signi"cant fraction of Li atoms from initially interstitial sites Li to substitutional sites Li due to a Coulomb-force-driven reaction Li #V NLi. We assume that Li in GaN occupies substitutional Ga-sites at temperatures above 4103C. A lattice site change was also observed for deuterium in GaN. Using nuclear-reaction analysis (NRA) channeling, Wampler et al. [25,27] observed H in the central region of the c-axis after ion implantation at room temperature. Annealing to 5513C and above causes the H to move from the interstitial sites to locations that could not be resolved with this technique. No deuterium remained at interstitial sites after annealing at 8093C. This temperature range is higher compared to that for the onset of Li di!usion, but close to the H redistribution temperatures reported by Wilson et al. [63,64,66,67] and the damage recovery reported above. Therefore, the lattice site change of Li can be unambiguously attributed to the di!usion of Li; whereas, the reaction of deuterium with vacancies or defects may be not only caused by deuterium di!usion, but also by di!usion of the defects during damage recovery. Further EC-investigations showed that sodium also occupies interstitial sites with no change upon annealing to 8003C [60]; whereas, the majority of the implanted heavy ions of In, Sr, Tm and Yb ions are substitutional after implantation with no change upon annealing as high as 12003C [35,60,96]. A high substitutional fraction was also observed for ion implanted Ca [97], Te [33], Er [24] and Hf [36] using RBS/C and particle induced X-ray emission (PIXE). All of these implanted heavy ions are substitutional; however, they have damaged surroundings and are slightly displaced

15 362 C. Ronning et al. / Physics Reports 351 (2001) 349}385 Fig. 7. Normalized emission yield of alpha particles emitted from implanted Li probe atoms measured at room temperature in GaN along the c-axis direction during Li implantation. The gray scale represents the normalized yield with respect to the yield for an o!-axis random direction. (From Ref. [95]) Fig. 8. Normalized alpha emission yield in the c-axis direction as a function of implantation temperature for undoped GaN. The change from channeling (yield '1) to blocking (yield (1) around 410$253C shows the lattice site change of a signi"cant fraction of Li from interstitial to substitutional sites. (After Ref. [95]) from the ideal substitutional position, potentially associated with defects [33,35,36,88,96}98]. Most of the heavier atoms: Ca [97], In [35], Pr(Ce) [98], Er [24], Tm [96], Yb [96] and Hf [36] exist substitutionally on the Ga lattice sites as opposed to N lattice sites. Unfortunately, the sub-lattice site of the donor Te, was not determined [33]. Annealing does not change the apparent substitutional fraction of the heavier atoms; however, it decreases the damage surrounding the implanted atoms and increases the fraction of atoms at ideal substitutional positions [33,88,96}98]. Silicon does not appear to be substitutional or interstitial direct after implantation [99,100]. The implanted Si atoms remains randomly distributed with only &20% substitutional after annealing to 10503C. However, upon annealing at 11003C there is a drastic change in the lattice location of the implanted Si with &100% of the implanted species located at substitutional Ga lattice sites [99,100]. MoK ssbauer spectroscopy was performed on GaN implanted with radioactive Cs, which decays to Sn [101]. It was indirectly determined that the Sn existed in several states in GaN including substitutional species on both Ga sites and N sites, as well as associated with defects and with O. Upon annealing the percentages of Sn occupying the various states in GaN changed. Speci"cally, the fraction on substitutional Ga sites increased at the expense of the fraction associated with defects [101].

16 C. Ronning et al. / Physics Reports 351 (2001) 349} Fig. 9. Photoluminescence spectra taken at 14 K of GaN implanted with various doses of Ge. Implantations were performed at a temperature of 77 K using an implantation energy of 300 kev. The inset shows an expanded view of the near band-edge emission along with its associated LO-phonon replicas for the unimplanted GaN. 3. Optical properties The optical properties of GaN can be best determined using photoluminescence (PL) and cathodoluminescence (CL) spectroscopy. The quality of the GaN sample as well as the optical activation of dopants can be identi"ed by the existence of speci"c PL-lines or -bands, by the intensity of PL-lines and by their width and shapes [2,5,7]. The "rst reported PL-study on implanted GaN was reported in 1976 by Pankove and Hutchby [58]. In this research spectra of 35 potential donor and acceptor species were recorded after ion implantation. The conclusions drawn at that time should today be used with care, because the quality of GaN samples have been improved by several orders of magnitude since that time. However, it had been shown that ion implantation resulted in nonluminescent GaN material implying the resulting implantation damage consists of many nonradiative recombination centers. This "nding is in agreement with all present studies and can easily be checked on an implanted GaN sample, which becomes nontransparent in the visible and near UV range after ion implantation. Thermal treatments are required to restore the luminescence, as reported subsequently in this section in the topic regarding the e!orts for optical donor, acceptor and rare earth activation in implanted GaN Defects Fig. 9(a) and its inset shows a PL spectrum of the unimplanted GaN sample for comparison. Several features can be observed including an intense near band-edge emission at 358 nm

17 364 C. Ronning et al. / Physics Reports 351 (2001) 349}385 (3.464 ev), commonly labeled I, that originates from recombinations of free excitons and/or excitons bound to shallow donors [102]. Two well-resolved, associated longitudinal optical (LO) phonon replicas at 367 nm (3.379 ev) and 377 nm (3.289 ev) for the I peak can also be observed, indicating the high quality of the unimplanted GaN. The position of the LO-phonon replicas results in a phonon energy of 85$5 mev for these samples, which is in agreement with values in the literature [103}105]. A weak transition at &381 nm (3.25 ev), ascribed to donor}acceptor pair (DAP) recombination [106], is observed, along with two LO-phonon replicas at 3.16 and 3.08 ev, as a shoulder on the second LO-phonon replica of the I emission. The identity of the shallow donor(s) is unknown, although both native defects and extrinsic impurities, including oxygen and silicon, have been suggested [78,107]. The identity of the acceptor is also unknown, although carbon and/or magnesium are likely candidates [108}110]. The feature observed at &550 nm (2.25 ev) is commonly referred to as the `yellowa emission band [111]. The origin of the yellow-band emission in GaN is still unknown. However, it is most likely that a variety of defects and deep level impurities causes this band and it has already been attributed [112}114] to deep acceptor levels that arise due to point defects such as Ga vacancies allowing a transition from shallow donor states to the deep acceptor state. The exact identities of both the shallow and deep defects are unknown with both native defects and extrinsic impurities possible. The broad, Gaussian line shape of several ω half-widths is typical of deep levels in semiconductors. The deep recombination partner has a large binding energy that results in a localized wavefunction (x) and thus a large range of k values [115]. Therefore, the transition occurs over a large energy range. As mentioned above, Fig. 9b}d clearly shows that the photoluminescence of GaN is severely depressed directly after ion implantation of Ge. The near-band-edge emission (3.464 ev) intensity is reduced by three orders of magnitude compared to the unimplanted sample even for a relatively low implantation dose of 1 10 cm. The intensity is further reduced at higher implantation doses and is totally quenched at an implantation dose of 5 10 cm. Ion implantation at higher substrate temperatures results in a reduction in the implantationinduced damage by dynamic annealing during implantation that is visible in the PL-spectra [56]. However, the reduction is minor compared to the large concentration of defects introduced during implantation, as attested by the fact the I line intensity is still four orders of magnitude lower than that of an unimplanted sample for implantation temperatures around 9003C [56]. Fig. 10 shows the PL-results of isochronal (1 h) annealing for Ge-implanted GaN with a dose of 1 10 cm. The"rst signi"cant change in the PL spectra can be observed after an annealing temperature of 6003C. A slight decrease of the weak defect related PL-band centered around 425 nm and a slight increase of the yellow band is visible, which indicates that some species of defects are annealed out. This temperature range is in agreement with the recovery observed with PAC (see Fig. 6) and can be attributed to the onset of defect/vacancy di!usion. The intensity of the near-band-edge I emission increases after annealing at each annealing step with the intensity increasing by about three orders of magnitude at 11003C. However, the intensity of the I emission is still one}two orders of magnitude lower than that observed in unimplanted GaN indicating that the recovery is still incomplete at this temperature, which is in agreement with the presented XRD, RBS, and PAC results. (An AlN capping layer was used to stabilize the GaN surface during the high temperature anneals.) The increase in the I intensity is also the result of the recovery of implantation damage by annihilation of point defects. This can be deduced from the fact that

18 C. Ronning et al. / Physics Reports 351 (2001) 349} Fig. 10. Photoluminescence spectra taken at 14 K of GaN implanted with Ge at a dose of 1 10 cm and annealed at various temperatures for 1 h. The implantation was conducted at room temperature with an implantation energy of 300 kev. The solid arrow is positioned at the Ge-related transition. unimplanted GaN samples do not exhibit any signi"cant changes under the same annealing conditions. The appearance of a new emission peak centered at &372 nm (3.33 ev) is observed after annealing at 11003C for 1 h Donor doping The new emission line at 3.33 ev in Fig. 10 is very likely due to a Ge related transition, as this peak was not observed in unimplanted GaN samples. A similar peak has been observed for GaN doped with Ge during growth [116] and by PL using radioactive Ge-doped GaN [117]. The transitional nature of this line is not known at this time. However, it should be a donor}band transition or a DAP transition with Ge as the donor. Assuming a donor}band transition, the energy level of the Ge donor would be 170 mev below the conduction band, assuming a band gap of 3.5 ev for GaN. This is not consistent with the reported [118}120] shallow nature of Ge. However, the transition may be also a DAP transition between the Ge donor and an unknown, potentially defect related, acceptor level. In this case the Ge donor level can be shallow depending on the energy level of the unknown acceptor. The emission of the 3.33 ev line is more intense for a lower implantation dose of 1 10 cm compared to 1 10 cm. This can be explained by the fact that the higher implantation dose result in a greater concentration of implantation-induced defects that are still present after the 11003C annealing step.

19 366 C. Ronning et al. / Physics Reports 351 (2001) 349}385 Silicon-implanted GaN annealed at 11003C for 1 h shows a strong DAP emission at 3.25 ev with the associated two phonon replicas at 3.16 and 3.08 ev. Other groups have observed similar strong DAP emission for GaN implanted with Si [17,21]. On the other hand, the shallow donor level (30 mev) of Si in GaN and its associated signature in the broadening of the I -line is well known from GaN doped with this species during growth [122,123]. Therefore, the enhancement in the DAP emission must be explained by the fact that the implanted silicon act as a shallow donor and result in increased DAP transitions with an unknown acceptor level. This behavior is also seen in GaN which is doped with Si during growth [123]. A new and sharp PL-line arising at nm (&3.40 ev) is also observed in silicon implanted GaN after thermal annealing. However, this peak is due to defects created during the implantation procedure, as this line is observed with varying intensities after implantation of Li, Be, Ge, In, Mg, Ca, and Er (see e.g. Figs. 11 and 12). It is likely that this emission is produced by nitrogen or gallium vacancies due to acceptor or donor bound excitons, because it can appear also in unimplanted GaN samples; however, it is dependent on the growth conditions [124] Acceptor doping Magnesium implanted and high-temperature-annealed GaN shows a strong DAP emission at 3.25 ev with two associated LO-phonon replicas at 3.16 and 3.08 ev (see Fig. 11). This emission is similar to the intrinsic DAP emission observed in undoped GaN (Fig. 9). However, the 3.25 ev emission, along with its LO-phonon replicas, is commonly seen in GaN samples doped during growth with Mg [125}141] and exhibiting p-type activation. The emission is attributed to a donor}acceptor pair transition (DAP) between a Mg acceptor and a shallow donor [134}137,140]. Therefore, it can be concluded that optical activation of the implanted Mg atoms has been realized. The donor of the DAP transition is unknown but is assumed to be the intrinsic donor in GaN. The Mg binding energy is estimated to be 235 mev assuming a band gap of 3.5 ev, a donor binding energy of 30 mev, and a value of 15 mev for coulomb interaction. This value is in good agreement with the accepted value of 225}250 mev. Fig. 11 shows also PL spectra of GaN implanted with di!erent doses of Mg and annealed at 12503C. With increasing implantation dose the Mg-related DAP transition decreases, and "nally the GaN samples implanted with a dose of 10 cm show no DAP-transitions. One would not expect such a behavior for increasing acceptor concentrations. Therefore, a high amount of residual defects, which strongly depends on the implantation dose, is still present in the annealed samples. This is supported by the fact that the behavior of the intensity of the yellow band is opposite to the behavior of the DAP transitions and therefore much higher in the GaN : Mg samples implanted with higher implantation doses (see Fig. 11). The PL spectrum of a GaN sample doped during MOCVD growth with 5 10 cm Mg is also contained in Fig. 11 [142]. The I -line of this samples is less intense compared to the emission from samples implanted with low doses. This is due to the higher intensity of the DAP-transitions. However, the di!erences between Mg-implanted and Mg-doped GaN are minor in this region of the spectrum. By contrast, the intensity of the yellow `defecta band shows a signi"cant di!erence: it is not visible in the Mg-doped GaN; whereas, the yellow band dominates the PL-spectra of the high-dose Mg-implanted GaN. Again, this demonstrates that after this high-temperature annealing

20 C. Ronning et al. / Physics Reports 351 (2001) 349} Fig. 11. Photoluminescence spectra measured at low temperature of GaN as a function of Mg implantation dose. (a) 10 cm, (b) 10 cm, and (c) 10 cm. The post implantation annealing was performed at 12503C for 30 min under vacuum and the implantation energy was set to 60 kev. The samples were protected with an epitaxial AlN-cap during annealing. For comparison, the PL-spectrum (d) was taken from a GaN-sample that was doped with 5 10 cm Mg during MOCVD-growth. procedure defects are still present in the material even when most of the implanted Mg ions are optical activated. A new PL-line at 3.36 ev appears in GaN implanted with 1 10 cm calcium ions and subsequently annealed at 12003C for 1 h, as shown in Fig. 12 (top). This emission was also observed, as a shoulder on the I emission peak, for higher Ca implantation doses [56]. The Ca binding energy was estimated to be 140 mev. This is close to the value of 169 mev reported by Zolper et al. [69,70,72] for the activation energy of implanted Ca in GaN as measured by temperature-dependent sheet carrier measurements. Thus, this emission peak is due to optically activated Ca. Beryllium is a promising candidate for shallow p-type doping in GaN given its calculated ionization energy of &0.06 ev [143]. We found a Be-related transition at 3.35 ev after ion implantation and annealing, as shown in Fig. 12 [144]. The Be-related PL-line is very weak compared to the Ca and Mg cases, which may be due to the possible formation of Be-defect complexes [145] or due to the occupation of non-doping interstitial sites [146]. We attributed the new line to band}acceptor (ea) recombinations. For this case, the ionization energy of Be acceptors can be calculated to be 150$10 mev. However, this is in contradiction to the theoretical value noted above. Dewsnip and co-workers [147] observed this line at ev in GaN samples doped with Be during growth. They calculated the ionization energy to be 90}100 mev due to the

21 368 C. Ronning et al. / Physics Reports 351 (2001) 349}385 Fig. 12. Photoluminescence spectra measured at low temperature of GaN implanted with a variety of species: (a) calcium, 10 cm, 300 kev, ¹ "12003C; (b) magnesium, 10 cm, 60 kev, ¹ "12503C; (c) beryllium, 10 cm, 100#200 kev, ¹ "9003C; (d) lithium, 10 cm, 30 kev, ¹ "8503C; (e) indium, 10 cm, 30 kev, ¹ "8503C; and (f) erbium, 10 cm, 60 kev, ¹ "8503C. assumption that this line is a donor-to-acceptor transition. Temperature-dependent PL are necessary to determine the nature of this Be-related transition and thus to calculate the exact activation energy of Be which can also be determined via electrical measurements. Zinc is another acceptor that has been extensively studied in GaN due to its reported high luminescence e$ciency. Implanted Zn has been reported to have a strong emission peak at 2.87 ev (&430 nm) in GaN [34,43,44,148}150], which is consistent with GaN "lms doped with Zn during growth. However, the intensity of the Zn emission in implanted "lms is typically less than in situ doped GaN [34]. Improvement of the Zn emission intensity after ion implantation was observed by Strite et al. [43,44] by annealing the implanted GaN at high temperatures or for long periods of time using a high N overpressure to stabilize the GaN during annealing. Suski et al. [150] showed that the optical activation of implanted Zn increases with increasing annealing temperatures to 13503C under high N overpressure and to 15503C under high N #Zn overpressure. This provides more evidence that the damage induced by implantation is detrimental to the optical properties of GaN and shows that the implantation induced damage is still being annealed out at temperatures in the range of 1350}15503C.