III-V and III-Nitride engineered heterostructures: wafer …The transferred layer is nearly...

III-V and III-Nitride engineered heterostructures: wafer bonding, ion slicing and more

O. Moutanabbir a, S. Christiansen a, S. Senz a, R. Scholz a, M. Petzold b,

and U. Gösele a

a Max Planck Institute of Microstructure Physics, Weinberg 2, Halle (Saale), D 06120 Germany

b Fraunhofer Institute for Mechanics of Materials, Halle, Germany

Wafer bonding in combination with ion slicing emerges as a

flexible technology by which bulk quality thin layers can be transferred onto different host materials. In the first part of this paper a successful application of this process to transfer 4-inch InP layer onto Si wafer is demonstrated. The use of an SiO2 interlayer makes the fabricated heterostructures compatible with high temperature processes. In the second section we address the applicability of this process for 2-inch freestanding GaN wafers which exhibit a strong post-implantation bowing making any bonding exceedingly difficult. We describe the origin of this bow enhancement and present a novel strategy to manipulate it. Based on our approach a successful bonding of 2-inch free-standing GaN onto sapphire handle wafers is achieved. Finally, by using a variety of experimental techniques, we explore the atomic processes and structural transformations involved in H ion-induced GaN splitting. A plausible mechanistic picture is presented.

Introduction

The first demonstration in the mid 90’s of silicon thin layer transfer using ion-cut

drew a great deal of attention (1). This process presents an effective approach for the integration of bulk quality thin layers onto different substrates achieving a wide variety of heterostructures frequently unattainable by epitaxy. Moreover, the concept of repeated transfer from the same wafer makes the process economically very attractive. Figure 1 illustrates, very schematically, the scenario leading to the subsurface layer transfer. Ion slicing consists of hydrogen (H) and/or helium (He) ion implantation into a donor wafer before bonding it to a handle wafer. The physical and chemical interactions of the implanted species with radiation damage and their thermal evolution act as an atomic scalpel leading to the formation of extended internal surfaces and to the complete exfoliation of the top layer of the implanted wafer. This phenomenon can be easily controlled by adjusting the implantation conditions. However, the most critical step in ion-slicing has to do with wafer bonding. In fact, achieving high quality and thermally stable bonding requires atomically flat surfaces and the absence of any irregularities in the wafers to be bonded such as the long range waviness and the bow (2). In addition, the surfaces have to be adequately activated before bringing them into contact. This surface engineering step depends sensitively on the nature of the materials to be bonded and on the subsequent thermal treatments.

ECS Transactions, 16 (8) 251-262 (2008)10.1149/1.2982876 © The Electrochemical Society

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In this paper, we focus on ion-slicing of 4-inch InP and 2-inch freestanding (fs-) GaN wafers. In the first section, we will demonstrate how epitaxy-compatible InP-on-Si substrates can be achieved. In the second section, we will address different issues involved in the case of 2-inch fs-GaN. Finally, by using a variety of experimental techniques, we explore the atomic processes and structural transformations involved in hydrogen ion-induced GaN splitting. Understanding these fundamental aspects is vital in order to control and optimize the ion-cut process.

Figure 1: Schematic illustration of the two steps involved in the ion-cutting process.

4-inch InP-on-Si heterostructures for high temperature processing

InP-on-Si is a very attractive heterostructure making conceivable the monolithic

integration of electronic and optoelectronic devices in the same platform and the realization of all Si-based optical communications (3). In addition, these heterostructures can also be used as templates for the fabrication of cost-effective and high performance solar cells (4). However, due to the high lattice mismatch between InP and Si (8.1%), direct heteroepitaxy leads to large threading and misfit dislocation densities in the grown layers and deterioration of the optical and electronic properties. Wafer bonding, in combination with ion-slicing, provides an economically attractive approach to fabricate bulk quality InP-on-Si substrates. However, due to their rather different thermal expansion coefficients (Si ~ 2.6×10-6 K-1 and InP ~ 4.6×10-6 K-1), the realization and use of such heterogeneous substrates face thermal stress problems which usually lead to debonding before splitting.

The transfer of 2-inch and 3-inch films onto Si wafers was already demonstrated (5-7).

Extrapolating this concept to 4-inch InP wafers is technologically highly relevant. However, with a thickness of more than 600 µm, the interfacial thermal stress will be much more pronounced in this case. Innovative approaches are needed to circumvent this problem. A first successful attempt to transfer 4-inch InP layers was achieved by Singh et al. (8). In that work the InP wafers were coated with a 150-nm-thick spin-on-glass (SOG) layer before bonding to thermally oxidized Si(001) handle wafers. Using a SOG intermediate layer leads to a very high surface energy which helps in avoiding thermal stress-induced debonding (9). However, the advantage of SOG is diminished (if not canceled) by the fact that the fabricated heterostructures cannot sustain high temperature

ECS Transactions, 16 (8) 251-262 (2008)


in subsequent epitaxy and device fabrication. SiO2 grown by plasma-enhanced chemical vapor deposition (PECVD) would be the appropriate interlayer for high temperature processing of InP-on-Si substrates. However, due to relatively lower surface energy in the case of hydrophilic bonding, thermal stress problems will have a greater impact. Here, we propose an effective preannealing procedure to overcome these problems and achieve a successful InP layer transfer.

In this work, 4-inch semi-insulting (100) InP wafers (Freiberger Compound Materials

GmbH) were used. A 150-nm-thick SiO2 layer was deposited on InP by PECVD. To prevent undesired out-gassing from the PECVD oxide layer during subsequent heat treatments of the bonded wafer pairs, the wafers were annealed at 850 oC in N2 atmosphere after SiO2-layer deposition. Oxide-deposited wafers were subsequently mirror-polished using chemo-mechanical polishing (CMP). The surface RMS roughness was in the order of 0.2 nm/10×10 µm2 after CMP, ensuring the short range flatness required for direct bonding. The wafers were then subject to He+ ion implantation at -15 oC at 100 keV with a fluence of 5×1016 cm-2. The implanted wafers were bonded at room temperature to thermally oxidized Si(001) handle wafers using a Süss Microtec CL200 cleaner. The bonding process was carried out in a class 1 cleanroom. The quality of the bonded interface was confirmed using infrared transmission imaging set-up.

Figure 2. XTEM micrograph of InP layer transferred onto Si wafer using a PECVD

grown SiO2 interlayer. In order to strengthen the bonding, the bonded pairs were annealed at 150 oC for 24

hours. Unlike InP/SOG/Si pairs (8), however, annealing at 200 oC necessary to initiate the splitting process leads to the cracking and breakage of the bonded wafers before the layer transfer occurs. As mentioned above, this is a consequence of the thermal mismatch and the relatively weak surface energy in the hydrophilic bonding. Fortunately enough, our systematic studies of the influence of pressure on bonding stability revealed that the thermal stress can be easily circumvented by applying a pressure of 2 MPa during annealing at 200 oC. After annealing for 30 min, the pressure was released and InP splitting takes place. Figure 2 displays a cross-sectional transmission electron microscope (XTEM) image of the as-split InP thin layer onto Si. The transferred layer is nearly 550-nm-thick. The top 220 nm is heavily damaged. Atomic force microscopy (AFM) analysis shows that the surface is very rough [Fig. 3 (a)]. Obviously, any growth on such substrates will lead to poor quality device structures. Therefore, it is necessary to remove the residual defective layer. For this reason, the as-split InP-on-Si wafers were subject to a chemical mechanical polishing (CMP). To optimize this process, we explored several slurries and polishing parameters. The optimal procedure consists of using a mixture of syton and water (1:20) and applying a pressure of 3 kPa to the wafer jig. After polishing, the obtained surface RMS roughness was in the order of 0.5 nm/20×20 µm2 [Fig. 3(b)].

InP

SiO2 (PECVD)

Si

Surface

Damage layer

150 nm



The crystalline quality of InP layer was confirmed by x-ray diffraction (XRD). Work is underway to grow device quality heterostructures on our InP-on-Si substrates.

Figure 3. AFM images of InP-on-Si substrates immediately after splitting (a) and

after polishing (b). Note the difference in X-Y scales. The corresponding RMS roughness is indicated.

Bonding of H-implanted 2-inch fs-GaN

GaN and related materials hold promise of future solid-state lightening, advanced

optoelectronics and even high-power radio-frequency electronics (10). However, the potential of these devices is still limited by the quality of the grown heterostructures. Presently, GaN used in device fabrication is epitaxially grown on sapphire despite its poor lattice and thermal match to GaN. The densities of misfit and threading dislocations in GaN layers deposited on sapphire range typically from 108 to 1010 cm−2 whereby the efficiency of GaN devices is limited. High quality GaN bulk substrates can be produced by hydride vapor phase epitaxy growth of thick GaN layers on sapphire and subsequent separation from the sapphire substrate. The current cost of these freestanding wafers is still so high that the concept of transfer of many layers from one fs-GaN wafer to appropriate host substrates by ion-cut is technologically and economically highly attractive. Here, we address some issues related to the application of H-ion cutting of 2-inch fs-GaN wafers.

Achieving high quality wafer bonding presents the most critical step in the process

(11). In fact, one of the major problems faced in bonding of 2-inch free standing GaN wafers is the strong enhancement of the bow due to hydrogen implantation (12). In this work, we present a novel approach to manipulate the implantation induced bowing phenomenon. By strain engineering at the back-side of the H-implanted fs-GaN wafer, we achieved a bow reduction in the first place and consequently high quality direct wafer bonding of the H-implanted 2-inch fs-GaN wafers to sapphire handle wafers. This is a crucial step towards fs-GaN thin layer transfer onto foreign substrates.

Experimental details

~300 µm-thick 2-inch double side polished fs-GaN wafers were used in this study.

Some wafers were subject to room temperature hydrogen ion implantation under the optimal conditions of ion-cutting at the energy of 50 keV with a fluence of 2.6 × 1017 atom/cm2. Elastic recoil detection (ERD) analysis shows that the H-implantation depth is ~450 nm with a concentration peak of ~13% around 320 nm under the selected

RMS = 39.63 nm

RMS = 0.5 nm

(a) (b)



implantation conditions (13). The bow of the wafers was measured using a long range DEKTAK 8 stylus profilometer on a length of 4.5 cm with a horizontal resolution of 3 µm. Microstructural information about H-ion implantation induced damage was obtained using (XTEM) and (XRD).

For bonding experiments, a 100 nm-thick SiO2 layer was deposited on fs-GaN by

PECVD. SiO2-deposited fs-GaN wafers were annealed at 850 oC in N2 atmosphere to

avoid any undesirable outgassing. The polished fs-GaN wafers with SiO2-layer and sapphire wafers (handle wafers) received a 15 min piranha (H2O2:H2SO4 = 1:3) solution cleaning followed by a deionized water rinse and N2 gun dry. After these cleaning steps the surfaces were terminated with hydroxyl (OH-) groups necessary to initiate the bonding (2).

Post-implantation bowing

During the implantation of energetic H ions several physical and chemical processes

take place leading to a variety of radiation damage-related structures including interstitials, vacancies, hydrogen-point defect complexes, voids, and free hydrogen. XTEM images of as-implanted fs-GaN wafers show a broad damage band extending over a 300 nm-thick layer below the surface [Fig. 4(a)].

Figure 4. (a) XTEM image of damage band induced in GaN by H ion implantation at

50 keV with a fluence of 2.6 × 1017 atom/cm2. (b) X-ray θ/2θ scans of (0002) fs-GaN after H implantation.

Typical x-ray diffraction data of fs-GaN before and after H implantation are presented

in Fig. 4(b). We notice that H implantation creates a significant out-of-plane tensile strain which is detectable as interference fringes extending from the left side of the (0002) GaN diffraction peak. This out-of-plane tensile strain is accompanying an in-plane compressive stress. As it will be addressed below, this radiation damage-induced in-plane compressive-strained layer has undesirable consequences on direct bonding of H-implanted fs-GaN wafers.

Figures 5 display the profilometer measurements at the N face of the fs-GaN wafers

for three different 2-inch wafers labeled A, B, and C. We note that all three wafers are dome-shaped with an initial bow (dashed lines) in the order of 9.9, 21.8, and 23.75 µm for wafers A, B, and C, respectively.

Damage band

50 nm -4000 -3000 -2000 -1000 0

100101102103104105

Inte

nsity

θ/2θ [arcsec]

out-of-plane tensile strain

HRXRD (0002)(a) (b)



A spectacular increase in the bow is recorded after H-implantation under ion-cut conditions [Fig. 5, red/dark lines]. We found that post-implantation bow reaches 38.9, 59.7, and 62 µm for wafers A, B, and C, respectively. This strong enhancement is a result of the in-plane compressive strain induced by high dose H-implantation as evidenced by x-ray analysis [Fig. 4(b)]. The average stress in the damaged layer can be approximately estimated from bow measurements (12). By analogy to a heteroepitaxial compressively strained layer (14), the average stress in the damaged layer can be estimated using the modified Stoney’s formula (15):

f

s

tL

btE

2

2

31

∆⎟⎠⎞

⎜⎝⎛

−=

νσ [1]

where E is Young’s modulus of GaN, ν is Poisson’s ratio, ∆b is the variation of the wafer bow induced by H-implantation, L is half of the profilometer scan length, ts and tf represent the thicknesses of the substrate and the damaged layer, respectively. Using the data from Fig. 5A-5C, the average stress in the implanted region was found to be in the order of -1.4 GPa.

Figure 5. Bow measurements on three different 2-inch fs-GaN wafers labeled A, B,

and C: unimplanted (dashed lines), single side (red/dark lines) and double side H-implanted (green/light lines).

The bowing of H-implanted fs-GaN wafers strongly limits or even prohibits direct

bonding and, consequently, the application of the ion-cut process to transfer thin layers of high quality fs-GaN to other host materials. It is well known that for direct bonding the wafers should be as flat as possible (2). The tolerable bows depend sensitively on the material properties and wafer thickness. Unfortunately, the observed post-implantation bowing is too high to allow any contacting of the wafers to be paired and bonded. Indeed, our bonding experiments on as-implanted wafers show that the gap between the two

0 1 2 3 40

20

40

60 C

Scan distance [cm]

0

20

40

60

Bow

[µm

]

0

20

40

B

A



surfaces is too large to even permit van der Waals forces to come into play and to initiate the bonding process [Fig. 6(a)]. It is also worth mentioning that our bonding tests under applied forces (e.g. a load in the bonding machine) have led to a systematic breakage and cracking of the wafers to be bonded and therefore does not present a real alternative option [Fig. 6 (b)]. The use of an intermediate layer as adhesive to ‘glue’ H-implanted fs-GaN to the handle wafer can lead to room temperature bonding. However, this approach is not suitable in our case since the subsequent steps of our process involve thermal treatments at high temperatures, e.g. the splitting at ~700 oC and the heteroepitaxy of device structures on the transferred GaN layers at at least 1050 oC. To the best of our knowledge, the commonly used adhesive layers cannot stand such high temperatures. Therefore, to avoid any complication in the process, the bonding via adhesive layers was not considered in our case. For technological reasons, the main challenge in fs-GaN layer transfer remains the reduction of the post-implantation bow in order to meet the direct bonding criterion.

Figure 6. (a) Schematic illustration of the bow-induced gap at the interface of the

wafers to be bonded. (b) Optical image of H-implanted fs-GaN (a half of 2-inch wafer) bonded to 2-inch sapphire wafer annealed under a pressure of 200 kPa. Note that due to radiation damage fs-GaN wafer exhibit a change in color from transparent to golden-brown after H implantation.

Bow manipulation by strain engineering

As it is addressed above, the strain at or below the surface strongly influences the

wafer bow. Similarly, a compressively strained layer at the back side (the Ga-face) of the H-implanted fs-GaN wafer (implantation on the N-face of the wafer) can enhance the bow from the Ga face of the wafer, naturally, thereby reducing the bow at the N-face of the wafer. The final value of the bow depends directly on the amount of strain induced from the Ga-face. Our finite element calculations demonstrate that the bow can be effectively reduced by ~87% when the strain is locally manipulated at the back side of the implanted 2-inch fs-GaN wafers [Figure 7]. However, this will remain a simple theoretical approach which can be hardly achieved experimentally.

Figure 7. Finite element modeling of bow reduction by wafer back side stress

manipulation.

Bowed wafer

1.0% 0.6% 0.5%

Flat wafer

Implanted zone

Handle Bonded area

Cracks



A strained layer in the back side of H-implanted wafers can be induced by mechanical polishing, heteroepitaxy, plasma treatment, or radiation damage. Ion implantation appears to be the most evident choice in our case since it can be included without any additional setup in the ion-cut process. Ga-face H-implantation under the exact conditions will certainly equibalance the bow enhancement observed after H-implantation at the N-face and the final bow on the wafer surface to be bonded can be reduced down to its initial value. Consequently, direct bonding of these wafers will sensitively depend on the bow of the original wafers. From data reported in Fig. 5, we see that unimplanted fs-GaN wafers exhibit already a large bow which cannot be tolerated in the direct bonding. Hence, in this specific case, further decrease of the bow is required in order to meet the bonding conditions. By examining Eq. 1, we note that the bow variation ∆b scales linearly with the thickness of the damaged layer tf at a fixed value σ, and vise versa. This means that producing a larger damage band or a higher strain can decrease the bow below its initial value. Since the thickness of the damage band depends on the H-ion energy and strain value varies linearly with the implanted fluence, the bow can be manipulated by adjusting the implantation energy and/or fluence. In this study, we chose to increase slightly the implantation depth while keeping the same ion fluence of 2.6 × 1017 H/cm2. fs-GaN wafers were implanted from the back side at about 65 nm deeper than the first implantation on the N-face. Results of profilometer measurements after double side implantations are shown in Fig. 5 (green/light lines). Expectedly, all the wafers exhibit a bow slightly smaller than the initial value.

Bonding of H-implanted 2-inch fs-GaN for potential layer transfer

Using double-side implantation direct bonding of H-implanted 2-inch fs-GaN wafers

to sapphire wafers has become possible. Fig. 8 displays images of a bonded pair. This achievement presents a critical step towards heterointegration of high quality GaN thin layers, which will have a major impact in the various GaN based device technologies.

Although back-side implantation is found to be a very effective strategy to achieve

direct bonding, one may think that it will present an additional costly process step, which may hardly justify the need for ion-cutting to achieve cost-effective substrates at the industrial level. We are developing new approaches to overcome this potential limit (16).

Figure 8. Optical (left) and IR (right) micrographs of H-implanted fs-GaN bonded to

2-inch sapphire wafer. The few interface bubbles observed can be caused by several factors including: particles on the bonding surfaces, localized surface protrusions, and trapped air pockets.

2-inch



Mechanistic picture of H ion-induced fs-GaN splitting In this section, we address the underlying physics of ion-induced thin layer

exfoliation. The ion-cut process has been successfully applied to various semiconductor materials such as Ge, InP, and GaAs. Independently of the material, the interaction of the implanted species with the radiation damage seems to play the key role in the splitting process. Apart from Si which was intensively investigated (17), only few studies were devoted to investigate the atomic processes involved in the splitting of other semiconductors. GaN exfoliation induced by H implantation was first reported by Kucheyev et al. (18). However, detailed studies on the mechanisms of ion slicing of GaN are still missing. In addition to XTEM, ERD, and XRD, we used Rutherford backscattering spectrometry in channeling mode (RBS/C) and positron annihilation spectroscopy (PAS) to investigate H-defect interactions leading to the splitting process (13).

0 5 10 15 20 25 30

0.50

0.51

0.52

0.53

0.54

0 87 261 480 765 1127 1484

virgin

Depth [nm]

S-p

aram

eter

Positron energy [keV]

as-implanted

0 100 200 300 400 5000

5

10

15

20

25

30

0.0

0.2

0.4

0.6

0.8

1.0Hydrogen

Ato

mic

Dis

plac

emen

ts [%

]

Depth [nm]

displacement field

H C

once

ntra

tion

[x10

22 H

/cm

3 ]

Figure 9. (left) S parameter depth profile measured before (triangles) and after

(squares) H implantation. (right) H atom depth profile (circles) and implantation damage profile (red line) as deduced from ERD and ion channeling, respectively.

Several atomic processes take place during the implantation of energetic H ions

generating defects from both sublattices. Fig. 4(a) shows a broad damage band extending over a 300 nm-thick layer starting about 200 nm below the surface. No extended defects are observed at the implanted fluence. High magnification images of the implanted zone taken under focus (not shown) indicate the presence of a high density of nanoscopic bright spots of ~1-2 nm diameter. The change in their contrast during focus variation suggests that they are void-like structures. We name these nanoscopic voids nanobubbles. Fig. 9 (left) displays Doppler broadening S parameter depth profiles measured before and after implantation. In PAS, the line shape of the γ ray depends on the electron momentum distribution. In the regions distant from atomic nuclei, such as vacancies or voids, the line has a distinctly smaller high momentum tail. The S parameter measures the “sharpness” of the line, so it takes higher values in vacancies and cavities. In the implanted region an enhancement of the S parameter of 6.5 % is observed indicating the presence of open volume defects which support XTEM observations. Such an increase is too large to be caused by monovacancies; obviously small vacancy clusters are detected. Figure 9 (right) displays H and atomic displacement depth profiles in as-implanted GaN as deduced from ERD and ion channeling analyses, respectively. The data show that H distribution is peaked at ~320 nm in agreement with SRIM calculations (19). H concentration at the peak is found to be ~1.2 × 1022 H/cm3 corresponding to an atomic concentration of about



~13% which is ~3 times higher than the concentration needed for Si exfoliation (17). The atomic displacement field reaches a maximum of ~27% at a shallower depth around 285 nm. We estimate the displacements per ion to be ~2.5, much smaller than ~10 calculated at a lattice temperature of 0 K (i.e., without dynamic annealing) (19). This suggests that about 75% of Frenkel pairs recombine during the implantation process. This annihilation rate is relatively smaller than in the case of Si implanted under ion-cut conditions (17). Therefore, the dynamic annealing cannot explain the unusually high fluence required for the splitting of GaN. The nature of H-defect complexes and their thermal evolution may play the most critical role than the absolute amount of the surviving defects.

Figure 10. (left) RBS/C yields as a function of annealing temperature for GaN

substrates implanted with H at 2.6 × 1017 atom/cm2. (right) XTEM micrographs of H-implanted GaN annealed at different temperatures: 450 oC, 500 oC, and 600 oC.

The thermoevolution of H-induced damage is investigated next. In figure 10, we

summarize RBS/C and XTEM data obtained after annealing at different temperatures. We note that RBS/C yield slightly decreases after heating at 300 oC. This can be attributed to the return of some interstitial atoms to substitutional sites. RBS/C spectra recorded after annealing at higher temperatures show unexpected features. The first is a strong broadening of damage-related peak observed only at the left side for backscattered energies below 1.28 MeV. This asymmetric broadening appears to be related to H since it occurs around the region where H reaches its peak concentration [Fig. 9 (right)]. The second feature in the spectra is the increase of dechanneling beyond the implanted zone. Finally, annealing above 300 oC causes also a strong enhancement of dechanneling near the surface. Note that the increase of the peak intensity comes simply from the increase of dechanneling background. Interestingly, the dechanneling level decreases above 450 oC. In XTEM data, we note that annealing up to 450 oC does not trigger any significant morphological changes in the damage band. Similarly to the as-implanted sample, the implanted zone remains decorated with nanobubbles. A small increase in temperature above 450 oC leads to the formation of nanoscopic cracks or platelets parallel to the surface. Further increase in the annealing temperature induces large cracks leading to a complete exfoliation of a ~340 nm-thick layer. Our detailed XTEM data show that structural transitions from nanobubbles to platelets and from platelets to microcracks occur within temperature windows as narrow as 25 and 50 oC, respectively. 450 oC is

50 nm

450 oC

50 nm

500 oC

100 nm

600 oC

0.50 0.75 1.00 1.25 1.50 1.75

500 oC

600 oC

450 oC

300 oC

Random

Nor

mal

ized

Bac

ksca

ttere

d Y

ield

Backscattered Energy [MeV]

Virgin

as-implantedplatelets

microcrack



identified as the critical temperature at which transformation processes commence. Interestingly, at this temperature the dechanneling enhancement in RBS/C yield attains its maximum.

A possible origin of lattice distortion would be trapping of H2 molecules in

nanobubbles leading to the buildup of internal pressure. However, highly pressurized nanobubbles alone cannot explain the observed dechanneling since their stress field can hardly reach the surface. Since damage layer contains also other defect complexes such as self-interstitial clusters, one can suppose that their combined influence with hydrogen-induced internal pressure can increase the in-plane compressive strain causing a strong lattice distortion. This process will ultimately lead to a weakening of the atomic bonding. The system attains the criticality around 450 oC. The dechanneling decreases above this temperature suggestive of a partial relief of the internal strain following the formation of platelets parallel to the surface. These platelets define the fracture paths for the exfoliation. Interestingly, Doppler broadening measurements for samples annealed at T ≤ 450 oC were found to be identical to those of the as-implanted state. The absence of vacancy clustering in this temperature range indicates that the necessary voids for splitting assemble dynamically during the implantation process. This behavior differs completely from the evolution observed in Si where an important increase of void-like defects was found to precede the exfoliation (20). This remarkable difference in vacancylike defects thermal behavior observed between GaN and Si suggests that a general and predictive microscopic model of ion-cut process has to consider the intrinsic properties of the material, the nature of H-defect complexes, and point defects diffusivities. These aspects are still poorly understood for GaN. Additional systematic experimental studies and calculations would be highly valuable.

Conclusion

In summary, epitaxy-compatible InP-on-Si substrates were successfully fabricated by

thin layer transfer from 4-inch InP donor wafer. We also demonstrated that H-implanted 2-inch fs-GaN wafers can successfully be bonded to sapphire by the manipulation of wafer bow using back side implantation. This makes the heterointegration of bulk quality GaN layers possible by using wafer bonding and ion-slicing. By using a variety of experimental techniques, we investigated the critical structural transformations involved in splitting of GaN by H. We found that vacancy clustering during the implantation process leads to the assembly of 1-2 nm nanobubbles. Lattice distortion was observed around the implanted zone in the temperature range 300–450 oC. The distortions result from the internal pressure buildup inducing a strain field around damage region. Strain relaxes by way of Ga-N bonds breaking leading to the nucleation of platelet embryos of extended internal surfaces.

Acknowledgments The authors acknowledge contributions from F. Süßkraut and R. Krause-Rehberg

(Univ. of Halle-Wittenberg), M. Chicoine (Univ. of Montreal), O. Seitz and Y. J. Chabal (Univ. of Texas at Dallas). InP wafers were supplied by Freiberger Compound Materials GmbH. This work has been partly funded by German Federal Ministry of Education and Research (CrysGaN Project).



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III-V and III-Nitride engineered heterostructures: wafer …The transferred layer is nearly...

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