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Laser damage mechanisms of two conductive wide-bandgap semiconductor films - indium tin oxide (ITO) and silicon doped GaN (Si:GaN) were studied via microscopy, spectroscopy, photoluminescence (PL), and elemental analysis. Nanosecond laser pulse exposures with a laser photon energy (1.03 eV, 1064 nm) smaller than the conductive films bandgaps were applied and radically different film damage morphologies were produced. The laser damaged ITO film exhibited deterministic features of thermal degradation. In contrast, laser damage in the Si:GaN film resulted in highly localized eruptions originating at interfaces. For ITO, thermally driven damage was related to free carrier absorption and, for GaN, carbon complexes were proposed as potential damage precursors or markers.
© 2016 Optical Society of America
1. Introduction
Transparent conducting films in optoelectronic devices require both high optical transparency in the visible range up to the near infrared and high electrical conductivity for rapid switching and uniform electric field applications. A third requirement is emerging: transparent conductive films with higher optical strength to enable more compact, high power opto-electronic devices, including those found in high repetition-rated, high average power laser systems [1]. Due to their wide bandgap and high conductivity, degenerately doped metal oxides such as indium tin oxide (ITO) have been widely used as transparent conducting films [2]. In order to achieve higher laser optical damage performance while maintaining conductivity and transparency, it is of paramount importance to understand laser absorption and damage mechanisms so transparent conductive materials deposition parameters such as film thickness and doping levels as well as the underlying substrate can be rationally selected. Furthermore, understanding failure modes in optoelectronic devices subject to high laser fluences is useful, especially for repeated exposures involved in the routine operations of these devices. Here, single-pulse and multi-pulse laser exposure failures were investigated to understand the basic processes involved in laser induced damage and any accumulated effects [3].
Laser induced damage morphologies of sputtered ITO films with various pulse duration [4] and wavelengths [5] have been studied using microscopy, but mostly to determine damage ablation threshold for film removal and without consideration of the damage processes. Other studies have focused on laser induced damage in fluence regimes relevant to removal of ITO film for machining purposes rather than degradation of the film relevant to device operation [5,6]. The thickness-dependent optical damage has been investigated [7] but also with limited characterization of the fluence-dependent damage modes and their cause. We present here detailed features of laser “damage” revealed by changes in optical microscopy images and deduce possible film damage mechanisms and processes that are consistent with observations as a function of laser fluence and number of nanosecond (ns) pulse exposures.
Although ITO is one of the main transparent electrode materials in use today, we note that other widegap next-generation semiconductor materials are now commercially viable and exhibit superior thermomechanical properties that should improve their optical strength. For example, widegap GaN with high carrier mobility can be grown with doping, resulting in a thick film with low in-plane sheet resistance [<10 Ω/square(sq)] but with fewer free carriers than degenerately doped thin ITO films. This is desirable because it results in less light absorption in the visible range up to near infrared wavelengths. Therefore, our studies also focus on the optical damage mechanisms in doped GaN. Previous laser damage studies of GaN, mostly conducted in the femtosecond regime accompanying the non-linear energy absorption, have focused on laser ablation for the purposes of machining thin films [8,9]. Yet, to our knowledge, GaN damage behavior for ns pulses at the fundamental Nd:glass 1064 nm wavelength where the GaN is transparent has not been reported. Furthermore, comparison of ITO and GaN laser damage mechanisms should provide insights into the rational design of transparent conductive films. To facilitate the discussion in the report, we summarized our observation in Table 1 below. In short, in ITO films with a high free carrier density >1 × 1021 cm−3, laser absorption produces damage due to gross heating, beginning with darkened regions which correspond to thermomechanical stress-induced cracking of the film; for GaN films doped with a free carrier density of ~1 × 1019 cm−3, laser damage involves localized absorption and eruption events which occur at a higher fluence than the damage in ITO films. Both GaN and ITO films exhibit incubation phenomena that increase their tendency to damage upon repeated exposures, where damage occurs at a lower fluence compared to that of single pulse damage.
Table 1. Single shot laser damage characterization as a function of laser fluence.
2. Experimental methods
2.1 ITO and GaN samples
The ITO films used in this study (WTSQ11050-C) were obtained from Thorlabs. The 90 nm ITO films were deposited using an e-beam evaporator onto a 5 mm thick N-BK7 substrate (1” square) that has a near-infrared (NIR) anti-reflection coating layer on the exit side. The carrier concentration (Ne) was measured as 1.99 × 1021 cm−3, the mobility as 29.5 cm2/Vs, and sheet resistance as 11.8 Ω/sq using a Hall effect measurement instrument (Ecopia, HMS-3000). Conductive GaN film samples were obtained from MTI Corporation and Kyma technologies. The 5 μm thick, intentionally Si doped n + GaN films (Si:GaN) were grown on a 300 nm AlN buffer layer supported by an Al2O3 substrate by hydride vapor phase epitaxy (HVPE). The 2” round sapphire c-plane substrate was offcut 1.0° towards the a-plane, and its exit side was polished. The carrier concentration was measured as Ne = 2.18 × 1019 cm−3, the mobility as 74 cm2/Vs, and sheet resistance as 7.7 Ω/sq using the Hall effect measurement instrument with non-magnetic phosphor-bronze contact probe tips with gold coating.
2.2 Laser damage experiments
Laser damage experiments were conducted under ambient conditions using a Nd:YAG laser (Spectra Physics, Quanta-Ray Pro 350). The YAG laser used in this testing contains an unstable resonator optimized for high output energy which is then focused in the far-field to a Gaussian spatial profile. The FWHM pulse duration was 9 ns and the wavelength was 1064 nm. The 1/e2 beam FWHM diameter was 650 μm with Gaussian beam profile (DX: 580 μm, DY: 720 μm). The angle of incidence was ± 1° while the samples were held vertically with film facing the laser. For the single pulse laser damage experiments, a single laser pulse was applied at a pristine location on the film. For multiple pulse laser damage experiments, multiple pulses were applied at a single location with a repetition rate of 10 Hz. The power stability of the laser was estimated to be ~ ± 3% over extended use.
2.3 Damage characterization
Optical microscope images and confocal laser scanning microscope images were obtained using a 3D Laser Scanning Confocal Microscope (Keyence, VK-X100) with two illumination sources – a white light lamp and a 658 nm laser. In order to create a 1.6 mm × 1.6 mm stitched microscope image, a high speed 1.3 mega pixel CMOS imaging camera (MC1311, Mikrotron) with a 20 × objective lens was used with custom software written in LabVIEW. For the photoluminescence (PL) 2D map, PL imaging was performed using a 3.1 eV pulsed laser (LDH-P-C-405B, Picoquant) and an objective (Mitutoyo 20 × , NA: 0.42) [9]. Luminescence excited by the laser was collected by the same objective. A spectral channel (550 - 610 nm) was monitored by an avalanche photodiode (Micro Photon Devices PDM 50CT). Data acquisition and analysis and sample translation (GTS, Newport) were performed using custom software written in LabVIEW. During PL spectral measurements, the sample of interest was directly irradiated using a 351 nm UV laser (Crystalaser) at an angle of incidence of 30° with respect to the surface normal. A 20 × objective (Mitutoyo, NA: 0.42) located at normal incidence to the sample surface was used to collect luminescence spectra to a spectrometer (Symphony II, Horiba) with a 150 groove grating (iHR320, Horiba). For Raman and PL spectrum measurements, a luminosity standard (Model 63355, Oriel-Newport) was used for spectrum calibration. SEM images (Nova NanoSEM 650, FEI) were obtained without metal coating. X-ray diffraction (XRD) measurements were performed in a θ/2θ configuration using a Bruker D8 Advance powder diffractometer (Cu Kα source, λ = 1.5406 Å) to determine the phase of the thin films. The grain size of crystallites within the ITO film was estimated via analysis of the XRD peak broadening using the Scherrer equation:, where τ is the mean size of the crystalline domains, Κ is a shape factor (0.9), λ is the wavelength of the incident radiation, β is the line broadening at the full width half maximum, and θ is the Bragg angle. The magnitude of β was determined via fitting of the XRD peaks using a self-consistent, quasi Monte-Carlo, algorithm. A combination of two peaks, corresponding to features arising from the Cu Kα1 and Cu Kα2 radiation, was used to fit each ‘peak’ in the experimental data. The intensity and separation of the Cu Kα1 and Cu Kα2 peaks were constrained within the model to reside within physically viable limits.
3. Results and discussion
3.1 Indium tin oxide
Optical failure modes of the conductive ITO films (10 - 100 Ω/sq) on glass substrates (summarized in Table 1) were investigated with single laser pulses across a range of laser fluences from 0.5 up to 11.0 J/cm2 at a wavelength of 1064 nm and a pulse length (τp) of 9 ns. Figure 1 shows microscope images of an ITO film damaged under laser irradiation, revealing increased degradation of the film with increased laser fluence. The apparent damage, which we defined as an observable modification of the film appearance with an optical microscope, started with “darkening” of the film [Fig. 1(a) at 4.0 J/cm2] and with increasing fluence, eventually developed common features of thermal degradation such as thermal stress cracking, melting, evaporation, and ablation. In Fig. 1(a), the single shot damage sites at 5.0 and 6.0 J/cm2 showed all of the features of thermal material degradation, while fewer of thermal degradation features were observed at lower fluences (4.0 and 4.5 J/cm2). Material removal of the ITO film in the center, accompanied by raised rims from Marangoni driven melt flow [10], was visible in the damage depth profile [Fig. 1(a) at 6.0 J/cm2]. Damage was confined to the 90 nm thick ITO layer and did not penetrate into the substrate.
Fig. 1 Optical micrographs of a single shot damaged ITO sample at a range of laser fluences in bright-field mode. (a) Damaged sites irradiated at the fluence range (0.5 - 3.0 J/cm2) do not show apparent film surface modification. All images in (a) have the same scale bar of 100 μm. The damage depth profile across the damage site is illustrated on the image at 6.0 J/cm2. (b) White dashed-brackets indicate the extent of the apparent affected area used in the graph (c). The laser beam profile is illustrated on the image at 8.0 J/cm2. All images in (b) have the same scale bar of 300 μm. (c) The apparent affected area is defined based on observation with an optical microscope and is plotted as a function of fluence.
A more detailed view of the laser damaged site in the ITO film ([Fig. 1], 4.5 J/cm2) is shown in Fig. 2. A stitched SEM image displays locations from the outside edge of the damage site all the way to the center. The numbered dots shown in Fig. 2(a) and 2(b) indicate the locations of the images in Fig. 2(c), for the purposes of direct comparison between the corresponding optical and SEM images [Fig. 2(b)]. In order to explain these images, the laser beam profile was assumed to be Gaussian, which would result in a nearly (super) Gaussian surface temperature distribution [11]. Under this profile, the temperatures are near ambient at the edge of the beam, while the peak temperature reached at the center is in excess of the films evaporation temperature in some cases (>2000 K for ITO) [5]. This temperature profile simultaneously produced a range of thermally activated film transformations that were apparent in both the optical and SEM images. In Fig. 2(c) near the outside edge of the site (frame 1), the film began to crack. The “darkening” can now be attributed, at least in part, to the scattering of light from cracks in the ITO film. The ITO film morphology in between those cracks remained unaffected in comparison to pristine areas of the film. The polycrystalline structure of the film was confirmed by XRD [Fig. 3] measurements. The grain sizes observed in SEM images were ~20 nm, which was comparable to the crystal grain sizes of ~10 nm that were calculated from Scherrer analysis of the XRD peak broadening. The phenomenon of cracks following the contour of the grain boundaries [12,13] could be caused by the mismatched coefficients of thermal expansion (CTE) of the glass substrate (αs = 8.3 × 10−6/K) and the ITO film (αf = 6.7 × 10−6/K for In2O3 [14]), or the rapid heat induced strain during ns laser exposure. A simple calculation of the mismatched thermal stress between the film and the substrate (the Young Modulus E = 119 GPa [14], the Poisson ratio ν of 0.35 [15], and the peak temperature gradient taken as ΔT~1000 K) suggested a tensile stress up to σ≈E(αs - αf)ΔT/(1-ν) = 0.3 GPa, which is well below the ITO yield strength of 1.2 GPa at ambient conditions. Therefore, dynamic and temperature dependent material properties may play a role in facilitating crack formation. Another mode of film crack formation could result from rapid ITO volume change due to fast heating/cooling or/and film melting and re-solidification processes analogous to “crack formation in a drying lake bed”. We also note that Ga-doped ZnO (Ga:ZnO), another popular thin film transparent conductive oxide, exhibited similar thermal stress cracking damage morphologies [Fig. 4] as those shown here for ITO but at lower fluences due to the lower thermal stability of the ZnO film with respect to Zn out diffusion [16] and lower mechanical hardness [17]. Due to its much lower stability and similarity with ITO failure modes when exposed to near infrared laser irradiation, Ga:ZnO films were not considered further.
Fig. 2 (a) Stitched SEM image of single pulse laser damaged site of ITO film at 4.5 J/cm2. The scale bar is 10 μm. (b) The numbered dots in the optical microscopic image correspond to the dots in (a). The scale bar is 100 μm. (c) Magnified SEM images of the numbered regions. Images are in the same scale, where the scale bar is 400 nm.
Fig. 3 X-ray diffraction patterns recorded in a θ/2θ configuration for the three samples: ITO (green), Ga-doped ZnO (blue) and Si-doped GaN (red). Each XRD pattern is normalized to the highest intensity peak in the data. Diffraction peaks associated with the substrates beneath the samples are denoted by a star (*). The peaks in the diffraction patterns for the Si-doped GaN and the Ga-doped ZnO are consistent with wurtzite-based structures. The ITO data displays a broad feature centered at ~26.5°, that is attributed to the amorphous substrate. A series of peaks arising from the ITO is also observed in this data. The breadth of these peaks suggests that the crystalline phase is composed of nano-scale crystallites (quantification of the crystallite size from the peak FWHMs was attempted via Scherrer-type analysis, which is described in the main manuscript).
Fig. 4 (a) Optical microscopic image of ZnO film damaged at 0.75 J/cm2. Upon single laser pulse irradiation, the reflection color from the damaged region changed. The scale bar is 300 μm. (b) SEM image of the damage site in (a). The scale bar is 100 μm. (c) Magnified SEM image of the center region of (b). The scale bar is 4 μm. The 300 nm Ga-doped ZnO film (4.67 Ω/sq, 1.4 × 10−4 Ωcm) as deposited by pulsed laser deposition technique at Pacific Northwest National Laboratory. The same laser illumination parameters previously discussed in the manuscript was used for the ZnO film damage testing.
Closer to the center of the damage site in Fig. 2(c) (frames 4 - 7), the film surface morphology began to change considerably with grain growth. The grain sizes were observed to increase by a factor of ten and the surface also became smooth via melting and surface transport [18]. The changes along the radius of the damage site were observed in optical microscopic images [Fig. 1 and Fig. 2(b)] in the way light is reflected, scattered, and absorbed. In the regions [frames 8 - 10 of Fig. 2(b)], where film melting and surface evaporation was expected due to reaching higher temperatures, metal phase segregation is possible due to the thermodynamic meta-stability upon laser irradiation; therefore, metal rich nanoparticles could be observed as white dots. Based on the Ellingham diagram and the ITO phase diagram [19], In2O3 is more stable than SnO2 especially at high temperature, from which we propose that the white dots arise from Sn enriched nanoparticles. Also, evaporated material could be re-condensing on the surface with cooling, where the vapor phase becomes supersaturated and metal enriched particles can be deposited as oxygen gas is removed.
By increasing the fluence from 6.0 to 7.0 J/cm2 (τp = 9 ns), the size of the affected region due to the laser illumination grew larger than the beam footprint as illustrated in Fig. 1(b). The increase of the size (radius) of the affected region to greater than the beam size can be explained by hot plasma formation above the surface during the laser illumination. In the event that the coupled laser energy becomes larger than the ionization energy of a target material, a plasma could be formed. Inspection of [Fig. 1(c)] indicates that this plasma affected region grows larger with fluence, but stabilizes in size perhaps due to plasma shielding [20]. SEM images [Fig. 5(a) and 5(b)] reveal that the plasma affected zone extended well beyond the ablation pit and SEM images at higher magnification [Fig. 5(c) and 3(d)] exhibit the plasma effect occurring between 6 J/cm2 (no plasma) and 7 J/cm2, respectively. In Fig. 5(d), stronger SEM intensity throughout the plasma affected region indicates changes in the e-beam electrical response from the film surface due to redeposition of 10’s nm sized particles and aggregates. In optical images, this debris-covered area appeared darker due to scattering losses. Debris redeposition can be explained by gas dynamic effects from the plasma induced sideway expansion (blast wave theory) that was experimentally confirmed under ambient conditions [21–23]. However, nanoscale pitting and darkening was not detected in our film as previously found in laser-induced plasma scalded dielectric mirror coatings, likely because the pulse energy levels involved for those materials tend to be much much higher than for ITO [24].
Fig. 5 SEM images of the ITO sample regions tested at laser fluences of (a) 6.0 J/cm2 and (b) 7.0 J/cm2, where plasma induced modification is observed. (c, d) Magnified SEM images of the regions marked with a square and circle, respectively, in (b), indicating electrical property modification due to the plasma formation. The scale bars are (a, b) 500 μm and (c, d) 150 nm.
In order to see structural information such as cracks [25], 2D maps of PL intensity using a 405 nm pump source were measured in Fig. 6(a). For the damage site at a fluence of 4.0 J/cm2, i.e. when the film began to show apparent damage with the “darkening” shown in Fig. 1(a), the PL intensity was higher than background. We note that, over the laser fluence range (from 4.0 to 7.0 J/cm2), the “darkened” regions in Fig. 1(a) were registered to the region with high PL intensity. Magnified images in Fig. 6(b) with the PL intensity profile show that the high PL intensity region is associate with a network of cracks that were observed in SEM images in Fig. 2(c, frames 1-3). In contrast, for central regions of the damage sites with melting, grain growth, and crack reduction, as shown in Fig. 2(c, frames 9 - 10), the PL was below the background of the unexposed, pristine film region. At higher fluence (7 J/cm2), where film was removed and the underlying substrate was exposed, the PL progressively increased to well above background in regions because the BK7 glass substrate has broad PL signal due to the impurities (i.e., Na, B, Ba, K) [26].
Fig. 6 (a) 2D map of photoluminescence (PL) intensity of damaged ITO film at varied laser fluence, where PL intensity lineout across the damage center is present. (b) Microscope image (color) and 658 nm laser reflection intensity 2D map (gray) at laser fluence of 4.0 and 4.5 J/cm2, where the PL intensity lineout across the damage center is present. The scale bars are 10 μm.
Consequently, considering all the above data, three single pulse fluence regimes emerged with varying effects on ITO films [Table 1]. The first regime (lowest fluences used) produced no apparent damage based on observations with conventional microscopy or PL measurement. Still, subtle degradation of the film’s electrical, optical, and structural properties could occur that will be addressed in future research. In the second regime, beginning with the onset of apparent damage at higher intermediary fluences, thermally activated processes occurred, including stress cracking giving rise to apparent darkening of the film, melting and re-solidification and, finally, evaporation and removal of film. In the third regime, at the highest fluences used, the formation of a plasma induced affected zone extending well beyond the beam footprint was associated with re-deposition of ITO debris that appeared dark. These three fluence regimes are evident in the plot of the optically apparent affected size versus laser fluence in Fig. 1(c). The significance of the latter two damage modes for device applications using transparent conductive electrodes is that, under such extreme fluence conditions, catastrophic device failure is likely because film cracking and removal perturb both optical and electrical performance. At lower fluence when no apparent changes occur even within normal operating conditions, the impact on the transparent conductive film properties might be cumulative until failure, or significant degradation could occur over time, impacting the usable lifetime and performance of the film. Hence, we investigated the possibility of a gradual deterioration, such as a shift in the optical damage susceptibility, by performing repeated laser exposure at a specific location and monitoring changes in the ITO film.
Figure 7 shows microscope images of ITO film damaged under multiple laser pulse exposures at a fixed fluence of 3 J/cm2 which is below the single shot damage threshold as shown in Fig. 1. The number of exposures is indicated in the top left corner of the image. In the initial pulse irradiation in Fig. 7(a), film darkening was not observed; however, after the next exposure (pulse) at the same laser fluence, film darkening occurred, suggesting that some film transformation had occurred in the previous exposure, although un-detectable by microscopy (or PL). The resulting damage after two exposures at 3 J/cm2 was similar in appearance to the single pulse damage that occurred at 4.0 J/cm2 [Fig. 1(a)]. With additional exposures, the film exhibited increasing damage (morphology and radius), and the damage site diameter expanded as illustrated in Fig. 7(e). Thus, continued outward expansion of the damage site from repeated exposures indicates a cumulative film degradation process involving film weakening or increased absorption. Traces of this particular damage growth mechanism were readily observed in SEM images in Fig. 7(b). Nearly co-aligned parallel tracks of remnant film material were aligned along the edge of the damage site in concentric patterns. The tracks were separated by a distance of about 100 nm, which corresponds to the thermal diffusion length of ITO in the plane of the film, Lth. For a pulse length of τp = 9ns and thermal diffusivity, DITO, of 1.6 × 10−6 m2/sec, Lth ~2√DITOτp was 150 nm as illustrated in Fig. 7(c), which was comparable to the measured tracks separation of ~100 nm seen in the SEM images. Therefore, repeated laser exposures ratcheted up the outward growth of the damage site in steps of ~100 nm via an absorption and heat diffusion front where in the case of 1000 pulses, we estimate an increase in diamter of ~200 μm. The small pits near the damage site edge [Fig. 7(d)] were the manifestation of thin film instabilities, related to a dewetting of the film melt from thermally nucleated holes [27] also visibile in the SEM images. ITO film damage from repeated exposures was observed even with fluences down to 2.0 J/cm2 (data not shown), well below the single damage threshold of 3.69 J/cm2 [Fig. 1(a)].
Fig. 7 (a) Microscope images of ITO film damaged under multiple laser pulses at a fixed fluence of 3 J/cm2 (beam dia. is 650 µm). The scale bar is 200 μm. (b-d) SEM images of a 10 nm thick ITO film under 1000 pulses at various magnification. (b) The scale bar is 100 μm. (c) Magnified SEM images of the dashed area in (b). The scale bar is 10 μm. (d) Magnified SEM images of (c). The scale bar is 1 μm. (e) Illustration of damage growth upon multiple pulses.
The conductive ITO damage mechanims discussed above for single and repeated exposures clearly implicate linear absorption and thermally driven damage involving multiple damage modes, depending on the film temperatures reached. Such an absorption and optical damage mechanism is to be expected due to the high doping level of typical tranparent conductive oxides like ITO with carrier densities, Ne, near ~1 × 1021/cc. These carrier densities produce plasma edge absorption in the near infrared wavelengths close to the laser wavelength used in this study (1064 nm). From the Drude model, , where is the plasma wavelength, c is the speed of light, Ne is the carrier concentration, e is the electronic charge, ε∞ is the high-frequency permittivity, εo is the permittivity of free space, m* is the effective mass, the estimated plasma wavelength of the ITO film is ~890 nm (ε∞ = 4, m* = 0.35me, where me is the electron mass), which represents the wavelength above which free electron absorption increases dramatically [28,29]. Thus, 1064 nm wavelength laser energy can be absorbed by free carriers to drive thermally activated laser-induced damage in ITO films.
Therefore, in order to minimize near infrared laser absorption, we considered another transparent conductive film model system with lower Ne and potentially lower free carrier absorption, but with nominally similar transmission of 70 - 90% [Fig. 8] and sheet resistances (Rs); we use technologically relevant values on the order of 1 - 1000 Ω/sq. required for low RC time constant optical switching needs. To maintain high film conductance, the lower Ne can be compensated by having a thicker film in conjunction with a higher carrier mobilty, μe, since Rs is proportional to 1/(tf × μe × Ne), where tf is the film thickness.
Fig. 8 Measured transmission spectrum of ITO sample and GaN sample. The illumination source was irradiated on the input surface (the film side). For the ITO sample, the near infrared anti-reflection coating layer on the substrate backside causes the oscillation in visible wavelength range. The spectrum was measured using UV-visible scanning spectrophotometer (Shimadzu, UV-1601PC).
3.2 Silicon-Doped Gallium Nitride
For such a system, we selected intentionally Silicon-doped Gallium Nitride film (Si:GaN or just “GaN” in this study) with an Ne ~1 × 1019/cc (i.e., 2 orders of magnitude lower than the ITO tested) and the plasma wavelength (calculated for the doped GaN with m* = 0.21) is shifted to the far infrared (6 - 13 µm wavelength), instead of the degenerately doped ITO’s near infrared absorption (p ~1 µm). Since GaN has a wide bandgap of ~3.4 eV, like ITO (Eg ~4 eV), it is transparent down to UV wavelengths. Although the optical strength of ITO and GaN used in this study cannot be compared directly due to their different carrier densities and thicknesses, use of doped GaN as a transparent conductive material represents a reasonable approach to minimize near infrared laser absorption and optical damage. This approach consists of selecting materials with, ideally, a high mobility >100 cm2/Vs (compared to ITO typical μe of ~10 cm2/Vs) [30] that can be grown sufficiently thick. High carrier mobility in GaN is due in large part to its high crystallinity, low m*, and lower concentrations of dopant impurities that can act as scattering centers, although phonon scattering still affects mobility at room temperature. We addressed specific optical damage modes and mechanisms for GaN films with high film conductivity and transparency, that can be compared with those of the ITO films.
Figure 9 includes typical laser damage results for a 5 µm thick GaN film (7.7 Ω/sq) on a sapphire substrate with a sheet resistance lower than the ITO films (11.8 Ω/sq). The GaN film was exposed to single and multiple shots at fixed fluences and locations (1064 nm wavelength illumination). It becomes immediately apparent that 1) the optical damage tends to occur at higher fluences than ITO by a factor of ~2 for single shots and by a factor of ~5 for multiple shots, and 2) that the damage modes are substantially different than those of ITO. In the case of GaN, laser irradiation produced discrete and highly localized damage sites with much smaller feature size than the one observed in the ITO films [Fig. 1]. Further, the GaN laser damage was accompanied with eruptions of materials leaving mulitple hexagonal pit shapes. The pit shapes are related to the hexagonal crystal structure of wurtzite GaN and were all aligned because the film is a single crystal layer. At higher fluences, above the single shot “threshold” of ~7 J/cm2 (τp = 9ns), the laser-induced damage in the GaN film showed no sign of gross or bulk thermal degradation in contrast to the ITO laser damage sites that had large areas of film cracking, melting, and removal by evaporation. Instead, the number of damage sites for GaN tended to increase with fluence [Fig. 9(a)], but damage pits did not always register exactly with the on-axis peak fluence of the Gaussian beam, indicating that precursors susceptible to laser damage existed within the film. Hence, laser damage in GaN films was less deterministic than those in ITO films because the discrete damage precursors were not uniformly distibuted. Also, we noted that, for thin films, the standing wave electric field (SWEF) distribution from optical interferences due to film internal reflections plays a significant role in determining the local threshold fluence. In some samples tested, the lateral distribution of damage density matched the pattern of visible optical interference fringes at 1064 nm (data not shown). Film thickness variations of both AlN and GaN layers across the sample can be as large as ± 1 µm for the “5 µm thick” GaN films, causing interference fringes. Unlike ITO, no apparent plasma related effects were observed for single shot exposures up to ~20 J/cm2 fluences (data not shown), although for repeated laser exposures, plasma related damage effects were observed as dark reddish film surface features [Fig. 9(b)] for GaN films at fluence <10 J/cm2.
Fig. 9 Optical microscopic images of GaN sample damaged upon (a) single pulse at laser fluences of 7.0 J/cm2, 9.0 J/cm2, and 12.0 J/cm2 and (b) multiple pulses of 10, 100, 1000 shots at laser fluence of 7.0 J/cm2. The scale bars are 100 μm. (c) GaN damage morphology evolution with multiple pulses at laser fluence of 8.0 J/cm2 (1/e2 beam diameter of ~100 μm, τp = 5 ns). The scale bars are 20 μm.
Unlike ITO, the GaN film appeared unaffected under optical microscopy until right before the first damage observation [Fig. 9(c) refer to shots “0 - 9” and “10”] upon repeated exposures at single shot sub-damage “threshold” fluences. The damage then grew by widening of the pit [Fig. 9(c) shots “10 - 100”] upon multiple exposures that reduce the damage “threshold” for Si:GaN, as it did for ITO. The reasons for multi-shot incubation and weakening, and/or increased in absorption by GaN are not clear and will be addressed in future research. In previous femtosecond pulsed laser damage studies, GaN exhibited cumulative effects from irradiation with multiple laser pulses, even for photon energies below the bandgap [31,32]. However, for ultrafast pulses (intensity: 1013 W/cm2), non-linear absorption is usually the dominant energy absorption path. However, the ultrafast laser damage mechanism cannot be directly applied to our study regime with ns pulses (<109 W/cm2).
In the following section, we carefully consider the film structure, including depth profiling of the damage sites within the film to reveal any dependence on depth in the GaN film. Since the damage sites are localized, we determined the location (depth) of the damage initiation within the film using SEM and 3D confocal microscopy. Figure 10 provides high resolution images of typical damage sites in the GaN film. The as-grown GaN surface exhibited the step-flow surface morphology of typical GaN films grown on offcut sapphire substrates (c-plane with 1.00° offcut towards a-plane). Details of the transparent layered film structure are shown in Fig. 11(c) (inset) and Fig. 12, which consists of a multilayer structure composed of a ~3 µm Si:GaN layer, a ~2 µm GaN seed layer, and a ~0.30 µm thick Aluminum Nitride (AlN) buffer layer that resides on a sapphire substrate (430 µm thick). Figure 10(a) and Fig. 13 show SEM images of a typical GaN damage site (at 9.0 J/cm2) where an eruption exposed the material near the substrate at the bottom of the pit. Near the center of the damage site [Fig. 10(b)], molten gallium (Ga) rich metal ejecta were apparent as determined from energy-dispersive x-ray spectroscopy (EDS) elemental analysis [Fig. 14]. The Ga region marked with the red dot in the SEM image [Fig. 10(b)] and in the corresponding microscope image in Fig. 10(c) appeared bright under optical microscopy. At high temperatures, molten GaN is not thermodynamically favored at ambient pressures, resulting in a decomposition process instead. A Ga rich substance would arise from the GaN thermal decomposition product accompanied by liquid gallium droplet formation [33] from laser induced damage.
Fig. 10 SEM images of the GaN sample damaged at laser fluence of 9.0 J/cm2. (a) SEM image of the damage site with largest diameter of Fig. 9(a, 9.0 J/cm2). The scale bar is 20 μm. (b) Magnified SEM image of the area marked with rectangle in (a). The scale bar is 1 μm. (c) Optical micrograph of the damage site in (a) with focus at 4.485 μm deeper from the film surface. The red dots in (b) and (c) indicate the same location. (d) Optical micrograph and (e) SEM image of the two small adjacent damage sites of Fig. 9(a, 9.0 J/cm2). The scale bars of (c-e) are 20 μm. (f) SEM images of exposed damage sites (side view). The scale bars are 5 μm.
Fig. 11 (a-c) GaN laser damage sites initiated at 12.0 J/cm2. Confocal laser scanning microscope images of (a) 2D depth map and (b) 3D depth map. The scale bar is 50 μm (c) Damage depth distribution. The inset shows SEM cross section image. (d) The depth profile of two exposed damage sites at 9.0 J/cm2. The inset illustrates the location of black and red dots of the depth profile.
Fig. 13 High resolution SEM image of Fig. 9(a). The scale bar is 10 μm. The inset shows a confocal laser micrograph of the same site. The step-flow induced layered surface morphology was observed in the pristine region and the layered structure is revealed on the exposed edge and sidewall. The dark SEM intensity in the damage center indicates that the underlying substrate is exposed. The same region is presented in the shiny reflection in the inset.
Fig. 14 Energy-dispersive X-ray spectroscopy (EDS) microanalysis of GaN damage sites. (a) EDS spectra of pristine film region. (b) SEM images showing EDS measurement locations of GaN damage sites at 9.0 J/cm2. The weight ratios (Ga/N) of pristine region and the location A are 5.41 and 12, respectively. It suggests that additional gallium with respect to nitrogen was introduced during the laser damage process and the gallium is shown as shiny reflecting substance in the optical microscopic image of the inset. In location B, nitrogen was not detected from the EDS measurement, indicating that the underlying substrate is exposed as we observed in the SEM image of Fig. 10(a). (c) Cross sectional SEM image of laser damage sites. We intentionally created many damage sites on the GaN sample and cleaved the sample to get a cross sectional image, where we observed presumably gallium droplets by chance. The scale bar is 5 μm.
Mechanisms pertaining to the localization of the initial absorption and damage initiation event could be obtained by considering Fig. 10(d) and 7(e) together. These images displayed the same area by optical microscopy and SEM, respectively. The optical image [Fig. 10(d)] revealed a fully erupted damage site on the top right corner of the image, where a cone of the film was completely removed. The secondary site in the image (outlined by the brackets) indicated just deformation of the film without removal, since the film was still attached throughout, although it separated slightly from the rest of the film. That gap between the deformed material and surrounding film produced the interference fringes visible in Fig. 10(d), while the damage initiation spot exhibiting the optically reflective Ga rich phase was still apparent at the center. Deformation-only site was likely the result of insufficient absorption of laser energy into the film to cause removal of a piece (cone) of the film. In the SEM image of the secondary site [Fig. 10(e)], the damage was not apparent in the area outlined in bracket since no fringes or sub-layer features can be observed in the SEM image. The outer surface morphology was unaffected and the damage initiation was located beyond the depth sensitivity of SEM. Thus, the laser damage initiation spots originated from inside the GaN film structure rather than the outer surface. We propose that extreme heat led to a rapid (ns) solid-vapor phase change, resulting in rapid volume expansion and shockwave that would bulge or deform the film above the damage initiation site. Figure 10(f) shows a typical SEM cross sectional view of a GaN damage site that exhibits a complete eruption originating from the substrate interface (top) and from near the middle of the film. The reason for the particular localization of the initiation damage events becomes apparent when we analyze the depth profile of the damage pits and their depth distribution.
Figure 11(a) shows a typical 2D depth map of the GaN film damage sites using a confocal laser scanning microscope. The image captures the aligned hexagonal pit structure, as indicated by the white arrows. Figure 11(b) provides a 3D layout of the same structure. The corresponding depth profiles were very similar across all damage sites as illustrated in Fig. 11(d). Analysis of the inter-planar angles of the damage pit walls (relative to basal c-plane) indicates that they are predominatly inclined at 17°, 25°, and 32°, which corresponds (within 1°) to the inter-semipolar-planes 10-12, 10-14, 10-13 of wurtzite GaN, respectively. These particular planes can thus be interpreted as being the weakest “cleavage” type planes in GaN, although the existence of such weak bonding planes have not previously been reported for GaN. A careful analysis of the depth measurements of the damage pits based on data such as those shown in Fig. 11(a) and 8(b) revealed the distribution of localized absorption [Fig. 11(c)] within the 5 µm thick GaN film. In this interpretation, the bottom of the damage pits was the location of the laser absorption and damage initiation event. Three distinct laser damage initiation locations were thus identifiable, and all are co-located or correlated with the interfaces within the film layered structure. The first was the sapphire substrate/film interface from which ~60% of all damage sites erupted. The second was the interface between the two stacked GaN layers [labeled “GaN-1” and “GaN-2”, respectively in Fig. 11(c) inset and Fig. 12] representing ~30% of all damage sites measured. The third location was near the outer surface of the film exposed to air, representing ~10% of damage sites. The three damage pit depths is not surprising considering that impurities, point defects, and structural defects are well known to occur and co-locate in GaN interfaces grown on sapphire [34–36], even in the presence of an AlN buffer layer. Surfaces and interfaces typically contain a large number of recombination centers because of the abrupt termination of the crystal, which leaves a large number of electron states inside the bandgap. In addition, surfaces and interfaces were more likely to contain impurities since they were exposed to ambient conditions when transferred between reactors to build the GaN film layered structures.
To characterize the presence of such defects in the GaN film, PL spectral emission was measured using UV (351 nm) laser illumination in Fig. 15(a). Because the band gap (Eg = ~3.4 eV) of GaN was smaller than the UV pumping source (3.53 eV), the probing depth was very shallow < 100 nm. Thus, illuminating each side of the sample separately, ie., the GaN film side and the sapphire substrate side (Eg(sapphire) ~9 eV), provided the PL information of GaN coming from different depths, mostly from the GaN region near the outer surface and near the substrate/AlN/GaN film interface (Eg(AlN) ~6 eV). The absolute PL intensity between the two curves in Fig. 15(a) cannot be directly compared due to the different illumination and collection schematics used, but comparison of the relative peak intensities is possible within a given spectrum. The PL spectrum of GaN film near the outer surface showed a strong peak at 364 nm (3.406 eV) due to the near band edge emission (NBE) and the so-called “red” and “yellow” luminescence (RL and YL, respectively), but they were negligible [37] in Fig. 15(a). In contrast, the PL emissions from GaN region near the substrate/film interface, where the largest density of laser induced damage sites were observed [Fig. 11(c)], showed far less NBE emissions at 368 nm peak (3.369 eV) and a stronger broad PL band with a peak near 690 nm (1.8 eV). The broad band PL could be fitted by two peaks - the first at 676 nm RL (1.8318 eV with FWHM: 0.3358 eV) and the second YL at 591 nm peak (2.0946 eV with FWHM: 0.6838 eV) as indicated in Fig. 15(a). The RL and YL indicate higher defect densities of the GaN film near the interface. The relatively lower levels of NBE at the substrate/film interface could be explained by the presence of a larger number of dislocations that are non-radiative recombination centers [38], in contrast with the surface PL spectrum that exhibits larger NBE PL. Because the film/substrate interface has both an unusually high damage density and a relatively high YL [Fig. 15(a)], the PL emissions may thus be correlated to laser damage.
Fig. 15 (a) PL spectrum of GaN sample with different illumination directions. The green curve is obtained from GaN film surface and the red curve is from GaN film near the GaN/AlN interface. (b) PL and (c) Secondary ion mass spectrometry (SIMS) of two GaN samples with different damage performance. For SIMS, solid line and dot line were used for the GaN sample with φdam ~3 J/cm2 and φdam ~0.1 J/cm2, respectively.
The film surfaces and interfaces of the sapphire/AlN/GaN sample are both rich in PL centers [36] that can absorb sub-band gap light (< 3.4 eV), including the laser wavelength of 1064 nm (1.165 eV) used here. The defect related optical transitions upon laser excitation can couple heat into the matrix of the material, which can result in catastrophic optical damage from temperature driven feedback loops as reported extensively in sub-bandgap laser illuminated GaAs materials [39]. Yet, an experimental identification of specific defects for the sub-bandgap light absorption is difficult in doped GaN since the material contains many kinds of structural defects and impurities that can all act as sources of absorption and non-radiative recombination centers [36]. An attempt was made to identify laser damage related defects by selecting two GaN film samples grown under nominally the same conditions but which exhibited different damage levels. One GaN sample with a high damage “threshold” of φdam ~3 J/cm2 (τp = 3ns) was compared to another sample with a much lower damage “threshold” of φdam ~0.1 J/cm2. A “threshold” was defined as a single damage site occuring per cm2 using large area beam exposures. Both types of sample showed nearly the same transmission in VIS-NIR and electrical conductivity, along with the same thickness and layer structure. The PL spectra of these two types of material were obtained using a LED broadband UV illumination centered at 365 ± 20 nm, where the entire film depth could now be addressed since the light absortion depth at 365 nm is larger than the GaN film thickness for wavelengths > 375 nm. The spectra were obtained using the same UV illumination and collection conditions between samples so PL intensities and peaks could be compared directly. The resulting PL curves in Fig. 15(b) showed that the low optical damage performance film (φdam ~0.1 J/cm2) exhibited both larger YL levels (by a factor of 2 - 3) and lower NBE PL than those of the higher damage performance GaN sample (φdam ~3 J/cm2). Damage tests on a few other samples with strong YL also confirmed this correlations between laser damage performance and YL and NBE levels. A more complete characterization of these differences in PL and laser damage will be addressed in future studies.
To determine the role of impurities as a source of YL and potential absorption centers, each type of film above was analyzed by secondary ion mass spectrometry (SIMS) to measure the local elemental composition across the depth of the two GaN films above [Fig. 15(c)]. Both the oxygen (O) and carbon (C) impurities were of interest since they have recently been established as the main YL sources [40,41], especially from CN-ON complexes. In the SIMS data, the C content of the GaN sample with low damage performance (φdam ~0.1 J/cm2) was factor of ~2 larger compared to the one with high optical damage threshold GaN (φdam ~3 J/cm2). Since O concentration was in excess with respect to C concentration, the limiting factor for the formation of CN-ON complexes was the carbon concentration. Carbon and oxygen are very common contaminants in growth chambers. The GaN sample with low damage threshold had both stronger YL emissions and higher carbon concentration than the GaN sample with high damage threshold. Still, since impurites also tend to agglomorate and stabilize near microstructural defects, the YL maybe a marker of other absorbing structural defects rather than a source of laser absorption. Therefore, YL or structural defects such as dislocations that affect NBE PL, or both could be implicated in damage initiation, especially at interfaces where they are known to occur due to conditions inherent to GaN growth. Furthermore, these interfaces also include highly degenerate n-type regions, which could contribute to laser damage [42].
Figure 16 shows Raman spectrum of GaN sample The probe laser for the Raman system used a wavelength of 532 nm and a 100 × objective lens (Nikon, NA: 0.8). A liquid nitrogen cooled spectrometer (Symphony II, Horiba) with 1200 groove grating (iHR320, Horiba) was used. The spectrum (red) was obtained by focusing Raman laser to the substrate and three peaks located at 417 cm−1, 578 cm−1, and 751 cm−1 are attributed to vibration modes of crystalline alumina. In the Raman spectrum (black), two vibration modes located at 573.91 cm−1 and 666.07 cm−1 are attribute to E2 (high) modes of GaN and AlN, respectively [43–45]. The peak locations are red-shifted from its unstrained bulk modes at 567 cm−1 and 657 cm−1 in Fig. 15, respectively, indicating that the GaN film and the AlN buffer layer are under tensile stress from the substrate due to the thermal expansion mismatch and high temperature growth conditions (~1000 °C) [13,45,46]. The presence of Si doping relaxes stress in the GaN film, but increases dislocation density [47]. Based on the E2 peak shifts (7.71 cm−1) based on strain free E2 peak position of 566.2 cm−1, we estimated the biaxial film stress to be approximately 1.79 GPa (biaxial stress = peak shift/4.3 GPa) [48].
Fig. 16 Raman spectrum of GaN sample. The black spectrum is from the GaN surface and the red is from the interface as illustrated in the inset with dots.
4. Conclusions
Laser induced damage of ITO films and GaN films with a sheet resistance of ~10 Ω/sq were investigated upon ns laser pulse irradiation. Unlike ITO films, the sources of laser absorption in conductive transparent GaN films are highly localized, resulting in discrete film eruptions. This study suggests C-O complexes (or clusters) as possible absorbers stabilized near microstructural defects at interfaces that can be correlated to laser damage via YL PL emissions from GaN films. Yet, their microscopic origin and their roles in damage initiation still need to be resolved in more detail. Furthermore, the discrete and stochastic nature of laser damage in GaN indicates that large area optical damage testing is necessary to fully assess this material’s damage modes, which will be addressed in future studies. In contrast, the ITO film laser damage is driven by more deterministic free carrier absorption and bulk heating that results in thermally activated damage processes such as thermomechanical stress induced cracking, and film melting and evaporation. As suggested by the laser incubation behaviors of ITO and GaN films exposed to repeated exposures, the properties of both ITO and GaN transparent conductive films may still be affected, even without apparent laser damage. This can ultimately affect optoelectronic device performance exposed to repeated and extreme laser conditions.
Funding
U.S. Department of Energy (DOE) (DE-AC52-07NA27344); Lawrence Livermore National Laboratory (LLNL); Laboratory Directed Research and Development grant (15-ERD-057)
Acknowledgments
This work performed under the auspices of the U.S. Department of Energy by Lawrence Livermore National Laboratory under contract DE-AC52-07NA27344 within the LDRD program. The authors thank Scott Chambers and Tim Droubay for providing Ga:ZnO films for laser damage studies. The Ga:ZnO were produced at the Environmental Molecular Sciences Laboratory (EMSL), a National Scientific User Facility located at Pacific Northwest National Laboratory (PNNL) and supported by the U.S. DOE Office of Biological and Environmental Research. We thank Paul Ehrmann for transmission measurements.
References and links
1. W. S. Brocklesby, “Progress in high average power ultrafast lasers,” Eur. Phys. J. Spec. Top. 224(13), 2529–2543 (2015). [CrossRef]
2. S. Ishibashi, Y. Higuchi, Y. Ota, and K. Nakamura, “Low Resistivity Indium Tin Oxide Transparent Conductive Films. 2. Effect of Sputtering Voltage on Electrical Property of Films,” J. Vac. Sci. Technol. A 8(3), 1403–1406 (1990). [CrossRef]
3. S. Z. Xiao, E. L. Gurevich, and A. Ostendorf, “Incubation effect and its influence on laser patterning of ITO thin film,” Appl. Phys. Adv. Mater. 107, 333–338 (2012).
4. W. T. Pawlewicz, I. B. Mann, W. H. Lowdermilk, and D. Milam, “Laser-Damage-Resistant Transparent Conductive Indium Tin Oxide Coatings,” Appl. Phys. Lett. 34(3), 196–198 (1979). [CrossRef]
5. O. Yavas and M. Takai, “Effect of substrate absorption on the efficiency of laser patterning of indium tin oxide thin films,” J. Appl. Phys. 85(8), 4207–4212 (1999). [CrossRef]
6. D. A. Willis, “Thermal mechanisms of laser micromachining of indium tin oxide,” Photon. Process. Microelectron. Photon. III 5339, 313–320 (2004). [CrossRef]
7. H. Wang, Z. Huang, D. Zhang, F. Luo, L. Huang, Y. Li, Y. Luo, W. Wang, and X. Zhao, “Thickness effect on laser-induced-damage threshold of indium-tin oxide films at 1064 nm,” J. Appl. Phys. 110(11), 113111 (2011). [CrossRef]
8. S. Nakashima, K. Sugioka, T. Ito, H. Takai, and K. Midorikawa, “Fabrication of periodic nano-hole array on GaN surface by fs laser for improvement of extraction efficiency in blue LED,” Phys. Procedia 5, 203–211 (2010). [CrossRef]
9. T. Kim, H. S. Kim, M. Hetterich, D. Jones, J. M. Girkin, E. Bente, and M. D. Dawson, “Femtosecond laser machining of gallium nitride,” Mat. Sci. Eng. B-Solid 82(1-3), 262–264 (2001). [CrossRef]
10. J. H. Yoo, J. B. In, C. Zheng, I. Sakellari, R. N. Raman, M. J. Matthews, S. Elhadj, and C. P. Grigoropoulos, “Directed dewetting of amorphous silicon film by a donut-shaped laser pulse,” Nanotechnology 26(16), 165303 (2015). [CrossRef] [PubMed]
11. S. T. Yang, M. J. Matthews, S. Elhadj, V. G. Draggoo, and S. E. Bisson, “Thermal transport in CO2 laser irradiated fused silica: In situ measurements and analysis,” J. Appl. Phys. 106(10), 103106 (2009). [CrossRef]
12. M. Batzill and U. Diebold, “The surface and materials science of tin oxide,” Prog. Surf. Sci. 79(2-4), 47–154 (2005). [CrossRef]
13. L. L. L. Fong Kwong Yam, Sue Ann Oh, and Zainuriah Hassan, “Gallium Nitride: An Overview of Structural Defects,” in Optoelectronics - Materials and Techniques (Padmanabhan Predeep, 2011).
14. H. S. Lee, J. O. Bang, H. J. Lee, G. J. Lee, K. H. Chai, and S. B. Jung, “Analysis of Thermo-Mechanical Behavior of ITO Layer on PET Substrate,” in 2011 IEEE 61st Electronic Components and Technology Conference (ECTC) (2011), pp. 1796–1799.
15. D. G. Neerinck and T. J. Vink, “Depth profiling of thin ITO films by grazing incidence X-ray diffraction,” Thin Solid Films 278(1-2), 12–17 (1996). [CrossRef]
16. Y. Song, E. S. Kim, and A. Kapila, “Thermal-Stability of Sputter-Deposited Zno Thin-Films,” J. Electron. Mater. 24(2), 83–86 (1995). [CrossRef]
17. M. D. Drory, J. W. Ager, T. Suski, I. Grzegory, and S. Porowski, “Hardness and fracture toughness of bulk single crystal gallium nitride,” Appl. Phys. Lett. 69(26), 4044–4046 (1996). [CrossRef]
18. W. W. Mullins, “Flattening of a Nearly Plane Solid Surface Due to Capillarity,” J. Appl. Phys. 30(1), 77–83 (1959). [CrossRef]
19. W. J. Heward and D. J. Swenson, “Phase equilibria in the pseudo-binary In2O3-SnO2 system,” J. Mater. Sci. 42(17), 7135–7140 (2007). [CrossRef]
20. R. K. Singh, “Transient plasma shielding effects during pulsed laser ablation of materials,” J. Electron. Mater. 25(1), 125–129 (1996). [CrossRef]
21. G. Koren and U. P. Oppenheim, “Laser Ablation of Polymers in Pressurized Gas Ambients,” Appl. Phys. B-Photo 42, 41–43 (1987).
22. A. Miotello, R. Kelly, B. Braren, and C. E. Otis, “Novel Geometrical Effects Observed in Debris When Polymers Are Laser Sputtered,” Appl. Phys. Lett. 61(23), 2784–2786 (1992). [CrossRef]
23. R. N. Raman, S. G. Demos, N. Shen, E. Feigenbaum, R. A. Negres, S. Elhadj, A. M. Rubenchik, and M. J. Matthews, “Damage on fused silica optics caused by laser ablation of surface-bound microparticles,” Opt. Express 24(3), 2634–2647 (2016). [CrossRef] [PubMed]
24. F. Y. Genin and C. J. Stolz, “Morphologies of laser-induced damage in hafnia-silica multilayer mirror and polarizer coatings,” Proc. SPIE 2870, 439–448 (1996). [CrossRef]
25. P. E. Miller, J. D. Bude, T. I. Suratwala, N. Shen, T. A. Laurence, W. A. Steele, J. Menapace, M. D. Feit, and L. L. Wong, “Fracture-induced subbandgap absorption as a precursor to optical damage on fused silica surfaces,” Opt. Lett. 35(16), 2702–2704 (2010). [CrossRef] [PubMed]
26. K. Mishchik, A. Ferrer, A. R. de la Cruz, A. Mermillod-Blondin, C. Mauclair, Y. Ouerdane, A. Boukenter, J. Solis, and R. Stoian, “Photoinscription domains for ultrafast laser writing of refractive index changes in BK7 borosilicate crown optical glass,” Opt. Mater. Express 3(1), 67–85 (2013). [CrossRef]
27. R. Seemann, S. Herminghaus, and K. Jacobs, “Dewetting patterns and molecular forces: a reconciliation,” Phys. Rev. Lett. 86(24), 5534–5537 (2001). [CrossRef] [PubMed]
28. H. Köstlin, R. Jost, and W. Lems, “Optical and Electrical Properties of Doped In2o3 Films,” Phys. Status Solidi, A Appl. Res. 29(1), 87–93 (1975). [CrossRef]
29. M. Fox, Optical Properties of Solids (Oxford, 2001).
30. D. H. Zhang and H. L. Ma, “Scattering mechanisms of charge carriers in transparent conducting oxide films,” Appl. Phys. A-Matter 62, 487–492 (1996).
31. P. G. Eliseev, S. Juodkazis, T. Sugahara, H.-B. Sun, S. Matsuo, S. Sakai, and H. Misawa, “GaN surface ablation by femtosecond pulses: atomic force microscopy studies and accumulation effects,” Proc. SPIE 4065, 546–556 (2000). [CrossRef]
32. P. G. Eliseev, H. B. Sun, S. Juodkazis, T. Sugahara, S. Sakai, and H. Misawa, “Laser-induced damage threshold and surface processing of GaN at 400 nm wavelength,” Jap. J.Appl. Phys. Part 2-Letters 38(Part 2, No. 7B), L839–L841 (1999). [CrossRef]
33. D. D. Koleske, A. E. Wickenden, R. L. Henry, J. C. Culbertson, and M. E. Twigg, “GaN decomposition in H2 and N2 at MOVPE temperatures and pressures,” J. Cryst. Growth 223(4), 466–483 (2001). [CrossRef]
34. J. Elsner, R. Jones, M. I. Heggie, P. K. Sitch, M. Haugk, T. Frauenheim, S. Oberg, and P. R. Briddon, “Deep acceptors trapped at threading-edge dislocations in GaN,” Phys. Rev. B 58(19), 12571–12574 (1998). [CrossRef]
35. W. L. Wang, W. J. Yang, H. Y. Wang, and G. Q. Li, “Epitaxial growth of GaN films on unconventional oxide substrates,” J. Mater. Chem. C Mater. Opt. Electron. Devices 2(44), 9342–9358 (2014). [CrossRef]
36. M. A. Reshchikov and H. Morkoc, “Luminescence properties of defects in GaN,” J. Appl. Phys. 97(6), 061301 (2005). [CrossRef]
37. M. A. Reshchikov, A. Usikov, H. Helava, and Y. Makarov, “Fine structure of the red luminescence band in undoped GaN,” Appl. Phys. Lett. 104(3), 032103 (2014). [CrossRef]
38. S. Y. Karpov and Y. N. Makarov, “Dislocation effect on light emission efficiency in gallium nitride,” Appl. Phys. Lett. 81(25), 4721–4723 (2002). [CrossRef]
39. J. W. Tomm, M. Ziegler, M. Hempel, and T. Elsaesser, “Mechanisms and fast kinetics of the catastrophic optical damage (COD) in GaAs-based diode lasers,” Laser Photonics Rev. 5(3), 422–441 (2011). [CrossRef]
40. D. O. Demchenko, I. C. Diallo, and M. A. Reshchikov, “Yellow Luminescence of Gallium Nitride Generated by Carbon Defect Complexes,” Phys. Rev. Lett. 110(8), 087404 (2013). [CrossRef] [PubMed]
41. S. O. Kucheyev, M. Toth, M. R. Phillips, J. S. Williams, C. Jagadish, and G. Li, “Chemical origin of the yellow luminescence in GaN,” J. Appl. Phys. 91(9), 5867–5874 (2002). [CrossRef]
42. D. C. Look and R. J. Molnar, “Degenerate layer at GaN/sapphire interface: Influence on hall-effect measurements,” Appl. Phys. Lett. 70(25), 3377–3379 (1997). [CrossRef]
43. M. Kuball, J. M. Hayes, A. D. Prins, N. W. A. van Uden, D. J. Dunstan, Y. Shi, and J. H. Edgar, “Raman scattering studies on single-crystalline bulk AlN under high pressures,” Appl. Phys. Lett. 78(6), 724 (2001). [CrossRef]
44. M. Kuball, “Raman spectroscopy of GaN, AlGaN and AlN for process and growth monitoring/control,” Surf. Interface Anal. 31(10), 987–999 (2001). [CrossRef]
45. T. Kozawa, T. Kachi, H. Kano, H. Nagase, N. Koide, and K. Manabe, “Thermal-Stress in Gan Epitaxial Layers Grown on Sapphire Substrates,” J. Appl. Phys. 77(9), 4389–4392 (1995). [CrossRef]
46. M. Kuball, J. M. Hayes, A. D. Prins, N. W. A. van Uden, D. J. Dunstan, Y. Shi, and J. H. Edgar, “Raman scattering studies on single-crystalline bulk AlN under high pressures,” Appl. Phys. Lett. 78(6), 724–726 (2001). [CrossRef]
47. I. H. Lee, I. H. Choi, C. R. Lee, E. J. Shin, D. Kim, S. K. Noh, S. J. Son, K. Y. Lim, and H. J. Lee, “Stress relaxation in Si-doped GaN studied by Raman spectroscopy,” J. Appl. Phys. 83(11), 5787–5791 (1998). [CrossRef]
48. C. Kisielowski, J. Kruger, S. Ruvimov, T. Suski, J. W. Ager, E. Jones, Z. Liliental-Weber, M. Rubin, E. R. Weber, M. D. Bremser, and R. F. Davis, “Strain-related phenomena in GaN thin films,” Phys. Rev. B 54(24), 17745–17753 (1996). [CrossRef]
