LAST, Sapphire (Al2O3), LiAlO2 ,LiGaO2 Crystal Substrates

LSAT (Lanthanum Strontium Aluminum Tantalum oxide) crystal is a newly developed multi-functional substrate for high Tc super conductive film. It is grown by Czockralski method and good quality by accurate controlling of the temperature distribution and growth rate. LSAT crystal has less structure defect, no domain, no twin structure. It also has low dielectric loss in microwave band. Sapphire (Al2O3) crystal substrate has a combination of optical and physical properties that make it best choice for a variety of demanding applications.It has good thermal properties, excellent electrical and dielectric properties. We supply high quality and low cost Sapphire crystal substrate   The lattice parameter of Lithium Dioxogallate (LiGaO2) and Lithium Aluminium Oxide (LiAlO2) crystal can match with the Gallium Nitride film very well. The mismatch coefficients are 0.2% and 1.4% only for LiGaO2 and LiAlO2, much smaller than that of the common used substrates such as <0001> sapphire (14%), <100> MgO (3%), <0001> SiC (3.5%). The Gallium Nitride film is a very important material for blue, violet, UV and white LED. A substrate material that matches the film to be grown well is very important to get a nice GaN epitaxial film. Crystal Properties: Substrates LAST Sapphire (Al2O3) Crystal structure Cubic Hexagonal Melting point 2113 K 2040 oC Density (g/cm3) 6.74  3.98 Specific heat 0.57 J/g.k (@ 295 K) Thermal conductivity 5.1 W/m.k  25.12 (@100 oC) Thermal Exp.coe.(10-6 /k) 8.2  (@ 295 K) 5.8 Cell parameter (Å) a=b=5.468, c=7.729 a=4.748, c=12.97 Hardness (Mohs) 6.5 9 Orientation <100>,<110>,<111> Substrates LiAlO2 LiGaO2 Crystal structure mm2 orthorhomnic 422 tetragonal Melting point 1900 oC 1600 oC Density (g/cm3) 4.187 2.615 Specific heat 11.8 (@300 K) 7.1 (@393~973 K) Cell parameter (Å a = 5.1687 Å a = 5.402 Å Hardness (Mohs) 5 6.5 Orientation <100>,<110>,<111> <100>,<110>,<111> Last crystals typical sizes (mm): 15x15, 10x10, 10x5, 10x3  Sapphire (Al2O3) typical sizes (mm): 10x3, 10x5, 10x10, 15x15, 20x15, 20x20, ø15, ø20, ø1", ø2"  LiAlO2, LiGaO2 typical sizes (mm): 10x3,10x5,10x10,15x15, ø10, ø15 Thickness (mm): 0.5, 1.0  Fuente: http://www.toplent.com/last.htm Nombre: Franklin J. Quintero Asignatura: EES

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A GaInN/GaN/AlGaN photodetector structure

Structure of the layers.

Scheme of the structure

TEM image of the QW (made by J. Borysiuk)

The sample has been grown by Metal-Organic Chemical Vapour Deposition by P. Caban in Institute of Electronic Materials Technology (ITME), Warsaw, Poland

Interferences in the buffer layer.

The sample is very flat. So it is possible to observe nice interferences. Light is reflected from upper surface of the sample and from GaN/sapphire interface. The two beams are very well parallel so they interfere. From the interference pattern, we can determine width of the GaN layer (2.66 μm at the center).

Photograph of interferences made in white light.

Interferences in 532-nm, green light.

Absorption spectra of the sample.

Photoluminescence maps.

Measurements of photoluminescence give information about quality of GaN and GaInN layers. GaN-related peak is at 360 nm. Its intensity and wavelength are nearly constant on whole surface of the sample.

Photoluminescence spectrum. GaN-related peak is at 360 nm. Emission from GaInN QW is at 400 nm.

Map of intensity of the GaN peak.

Map of wavelength of the GaN peak.

GaInN QW related peaks have wavelengths between 390 and 420 nm depending on a position on the sample. the wavelength of emission from QW depends upon width of the well. The broader well - the longer wavelength. The circular shapes on the maps are due to substrate rotation during growth of the sample.


(A) Photograph of the sample. (B) Map of intensity of the GaInN QW peak. (C) Map of wavelength of the GaInN QW peak.

Photoluminescence has been measured by A. Kos and K. P. Korona in Institute of Experimental Physics, Warsaw University, Warsaw, Poland


Fuente: http://info.fuw.edu.pl/~kkorona/det_AlGaInN.html
Nombre: Franklin J. Quintero
Asignatura: EES

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Hot STM Studies of MBE growth of Si on Si(001)

We studied the initial stages of Si MBE growth on Si(001) surfaces. In the initial stages of MBE growth at 530 K, many small one-dimensional (1D) islands are formed. The explanation of this curious shape anisotropy has been controversial. On the Si(001)-(2 × 1) surface, the mobility of Si monomers and dimers is high along the dimer rows and low across them. It has been observed that for certain growth conditions, deposited Si atoms form islands a single dimer row wide (1D islands). The long axes of these 1D islands are aligned perpendicular to the substrate dimer rows. It is puzzling that the islands grow perpendicular to the direction of fast diffusion. Because the mobility is high along the rows, one might expect that islands would capture atoms from a greater distance in this direction. This would lead to growth predominantly along the rows, contrary to what is observed.

Mo et al. proposed a kinetic model based on sticking anisotropy to explain the island shape anisotropy. If island ends are stickier than island sides, arriving adatoms stick preferentially to the ends, resulting in the growth of many 1D islands. Tsao and co-workers suggested that an opening or weakening of dimer bonds might explain this sticking anisotropy. Metiu et al. proposed an alternative model based on an exchange mechanism. Metiu suggested that an adatom arriving on the side of a Si island may displace an existing island atom to the top of the island. The displaced atom diffuses rapidly along the top of the island (the direction of fast diffusion) until reaching an end, where it can fall over the edge and stick. Adatoms arriving on island sides are transported to the ends via this exchange mechanism, resulting in enhanced 1D growth. Our ability to track individual islands microscopically with the STM allowed us to test these models directly.

FIG 1. 400 × 400 Å STM Image of the Si(001) surface at 536 K showing many small Si islands. The direction of fast monomer diffusion is parallel to the dimer rows, indicated by the black arrows.

In Fig. 1 we show an image of the Si(001) surface decorated with many Si islands. Islands have grown on the upper and lower terraces with their long axes aligned perpendicular to the substrate dimer rows. According to the simple sticking anisotropy model of 1D island growth, a long 1D island should be no more likely to capture material at its ends than a short one, for both islands have just two ends. The Metiu model, however, implies that the effective end capture probability should increase linearly with length. Because the exchange mechanism transports adatoms arriving on island sides to island ends, doubling an island's length should double the rate at which material is added to its ends. A portion of one movie is shown in Fig. 2. From detailed analysis of many movie images, we found that growth is independent of length, supporting the anisotropic sticking model. One advantage of our method is that we can select and follow just the 1D islands. The simple anisotropic sticking model is therefore confirmed as the cause the island shape anisotropy observed during growth. We then use our data to measure the anisotropic sticking ratio, yielding a sticking anisotropy ratio of 0.019 ± 0.003. Thus, an end site is roughly 50 times more likely to gain a block than a side site.

Figure 2. 800 × 800 Å STM images from a longer movie showing the growth of 1D islands at 523 K. Time advances from left to right. Coverage increases from 0 to 0.07 ML. The movie begins with the clean substrate before deposition. One can follow individual islands from frame to frame as they grow. View Java Applet Movie or View GIF Animation

Figure 3. 800 × 800 Å STM images showing the growth of islands at 533 K. Time advances from left to right. Coverage increases from 0 to 0.1 ML at a rate of approximately 0.01 monolayers deposited per frame. The movie begins with the clean substrate before deposition. At this temperature 1-D islands form and coalesce. View Java Applet Movie or View Gif Animation

This result gives us insight into MBE growth of Si on Si(001). It is known that in step flow growth, type-B step edges grow faster than type-A edges, eventually causing double height steps to form. Although most material arriving at a step edge is incorporated at existing kink sites, the creation of new kinks (by addition of material to previously flat sections) is the rate-limiting step for the advance of the edge. Because the side of a 1D island is a type-A edge, adding a block there is like adding a block to a flat section of a type-A step edge. The end of a 1D island is a type-B edge; adding a block there is similar to adding a block to a flat section of type-B step edge. Thus the rapid growth of the type-B step edge and the highly anisotropic island shapes are both results of the sticking anisotropy.
Our measured sticking anisotropy, together with our previous measurements of edge fluctuations, also provides a detailed quantitative picture of the coarsening of small islands or features on the Si(001) surface. In this case, the type-B edges of each row fluctuate. The rows on the edge of an island fluctuate fastest, and when the ends cross, the entire row disappears (this process is very clear in Java Applet or Gif Animation movie images of island fluctuations) [Pearson 1995b]. Due to the sticking anisotropy, it is then difficult to nucleate a new row, and the island shrinks. This process will hasten the demise of smaller islands at the expense of larger islands and steps.
The actual microscopic mechanism underlying these step fluctuation and growth processes is not currently known. The current models include monomer, dimer, and dimer vacancy diffusion. The low prefactor that we observed for step fluctuations indicates that a multi-atom or collective process is involved. The activation energy of 0.97 eV that we measure is larger than the 0.67 eV estimated for monomer diffusion, and smaller than the dimer dissociation energy. On the other hand, 0.97 eV is very close to the roughly 1 eV activation energy for dimer diffusion. We suggest that for typical conditions, the monomers pair up into dimers.
Further investigation of epitaxial growth's early stages has therefore involved a detailed study of single and two dimer configurations and dynamics. This has lead to our recent observation of a novel dimer diffusion mechanism that provides for crossing substrate rows thereby extending the possible diffusion pathways into two dimensions. The Stealth, or C-type dimer configuration plays a key role in this mechanism. Surprisingly, the C dimer plays a similar role in the transformation between two four-atom configurations, a transformation which we have observed to lead to epitaxial row formation. Especially at elevated temperatures where process rates are increased, a conventional STM's image rate limits its applicability to surface dynamics studies. For this reason we have implemented a tracking technique which allows us to follow individual features along their diffusive paths. This has yielded dimer diffusion and dissociation activation energies.

Fuente: http://groups.physics.umn.edu/stmlab/growth/growth.html
Nombre: Franklin J. Quintero
Asignatura: EES

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Epitaxy

Epitaxy refers to the method of depositing a monocrystalline film on a monocrystalline substrate. The deposited film is denoted as epitaxial film or epitaxial layer. The term epitaxy comes from the Greek roots epi, meaning "above", and taxis, meaning "in ordered manner". It can be translated "to arrange upon".
Epitaxial films may be grown from gaseous or liquid precursors. Because the substrate acts as a seed crystal, the deposited film takes on a lattice structure and orientation identical to those of the substrate. This is different from other thin-film deposition methods which deposit polycrystalline or amorphous films, even on single-crystal substrates. If a film is deposited on a substrate of the same composition, the process is called homoepitaxy; otherwise it is called heteroepitaxy.
Homoepitaxy is a kind of epitaxy performed with only one material. In homoepitaxy, a crystalline film is grown on a substrate or film of the same material. This technology is used to grow a film which is more pure than the substrate and to fabricate layers having different doping levels. In academic literature, homoepitaxy is often abbreviated to "homoepi".
Heteroepitaxy is a kind of epitaxy performed with materials that are different from each other. In heteroepitaxy, a crystalline film grows on a crystalline substrate or film of a different material. This technology is often used to grow crystalline films of materials for which single crystals cannot otherwise be obtained and to fabricate integrated crystalline layers of different materials. Examples include gallium nitride (GaN) on sapphire or aluminium gallium indium phosphide (AlGaInP) on gallium arsenide (GaAs).
Heterotopotaxy is a process similar to heteroepitaxy except for the fact that thin film growth is not limited to two dimensional growth. Here the substrate is similar only in structure to the thin film material.
Epitaxy is used in silicon-based manufacturing processes for BJTs and modern CMOS, but it is particularly important for compound semiconductors such as gallium arsenide. Manufacturing issues include control of the amount and uniformity of the deposition's resistivity and thickness, the cleanliness and purity of the surface and the chamber atmosphere, the prevention of the typically much more highly doped substrate wafer's diffusion of dopant to the new layers, imperfections of the growth process, and protecting the surfaces during the manufacture and handling.
Applications

Epitaxy is used in nanotechnology and in semiconductor fabrication. Indeed, epitaxy is the only affordable method of high crystalline quality growth for many semiconductor materials, including technologically important materials as silicon-germanium, gallium nitride, gallium arsenide, indium phosphide and graphene.
Epitaxy is also used to grow layers of pre-doped silicon on the polished sides of silicon wafers, before they are processed into semiconductor devices. This is typical of power devices, such as those used in pacemakers, vending machine controllers, automobile computers, etc.

Methods
Epitaxial silicon is usually grown using vapor-phase epitaxy (VPE), a modification of chemical vapor deposition. Molecular-beam and liquid-phase epitaxy (MBE and LPE) are also used, mainly for compound semiconductors. Solid-phase epitaxy is used primarily for crystal-damage healing.

 Vapor-phase

Silicon is most commonly deposited from silicon tetrachloride in hydrogen at approximately 1200 °C:

SiCl_4(g) + 2H_2(g) ↔ Si_(s) + 4HCl_(g)

This reaction is reversible, and the growth rate depends strongly upon the proportion of the two source gases. Growth rates above 2 micrometres per minute produce polycrystalline silicon, and negative growth rates (etching) may occur if too much hydrogen chloride byproduct is present. (In fact, hydrogen chloride may be added intentionally to etch the wafer.) An additional etching reaction competes with the deposition reaction:

SiCl_4(g) + Si_(s) ↔ 2SiCl_2(g)

Silicon VPE may also use silane, dichlorosilane, and trichlorosilane source gases. For instance, the silane reaction occurs at 650 °C in this way:

SiH₄ → Si + 2H₂

This reaction does not inadvertently etch the wafer, and takes place at lower temperatures than deposition from silicon tetrachloride. However, it will form a polycrystalline film unless tightly controlled, and it allows oxidizing species that leak into the reactor to contaminate the epitaxial layer with unwanted compounds such as silicon dioxide.
VPE is sometimes classified by the chemistry of the source gases, such as hydride VPE and metalorganic VPE.

 Liquid-phase

Liquid phase epitaxy (LPE) is a method to grow semiconductor crystal layers from the melt on solid substrates. This happens at temperatures well below the melting point of the deposited semiconductor. The semiconductor is dissolved in the melt of another material. At conditions that are close to the equilibrium between dissolution and deposition the deposition of the semiconductor crystal on the substrate is slow and uniform. Typical deposition rates for monocrystalline films range from 0.1 to 1 μm/minute. The equilibrium conditions depend very much on the temperature and on the concentration of the dissolved semiconductor in the melt. The growth of the layer from the liquid phase can be controlled by a forced cooling of the melt. Impurity introduction can be strongly reduced. Doping can be achieved by the addition of dopants.
The method is mainly used for the growth of compound semiconductors. Very thin, uniform and high quality layers can be produced. A typical example for the liquid phase epitaxy method is the growth of ternery and quarternery III-V compounds on gallium arsenide (GaAs) substrates. As a solvent quite often gallium is used in this case. Another frequently used substrate is indium phosphide (InP). However also other substrates like glass or ceramic can be applied for special applications. To facilitate nucleation, and to avoid tension in the grown layer the thermal expansion coefficient of substrate and grown layer should be similar.

 Solid-phase

Solid Phase Epitaxy (SPE) is a transition between the amorphous and crystalline phases of a material. It is usually done by first depositing a film of amorphous material on a crystalline substrate. The substrate is then heated to crystallize the film. The single crystal substrate serves as a template for crystal growth. The annealing step used to recrystallize or heal silicon layers amorphized during ion implantation is also considered one type of Solid Phase Epitaxy. The Impurity segregation and redistribution at the growing crystal-amorphus layer interface during this process is used to incorporate low-solubility dopants in metals and Silicon.

Fuente: http://en.wikipedia.org/wiki/Epitaxy
Nombre: Franklin J. Quintero
Asignatura: EES

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Metalorganic vapour phase epitaxy

Metalorganic vapour phase epitaxy (MOVPE) is a chemical vapour deposition method of epitaxial growth of materials, especially compound semiconductors from the surface reaction of organic compounds or metalorganics and metal hydrides containing the required chemical elements. For example, indium phosphide could be grown in a reactor on a substrate by introducing Trimethylindium ((CH₃)₃In) and phosphine (PH₃). Alternative names for this process include organometallic vapour phase epitaxy (OMVPE), metalorganic chemical vapour deposition (MOCVD) and organometallic chemical vapour deposition (OMCVD). Formation of the epitaxial layer occurs by final pyrolysis of the constituent chemicals at the substrate surface. In contrast to molecular beam epitaxy (MBE) the growth of crystals is by chemical reaction and not physical deposition. This takes place not in a vacuum, but from the gas phase at moderate pressures (2 to 100 kPa). As such this technique is preferred for the formation of devices incorporating thermodynamically metastable alloys. It has become the dominant process for the manufacture of laser diodes, solar cells, and LEDs.

 Reactor components
A reactor is a chamber made of a material that does not react with the chemicals being used. It must also withstand high temperatures. This chamber is composed by reactor walls, liner, a susceptor, gas injection units, and temperature control units. Usually, the reactor walls are made from stainless steel or quartz. To prevent over heating, cooling water must be flowing through the channels within the reactor walls. Ceramic or special glasses, such as quartz, are often used as the liner in the reactor chamber between the reactor wall and the susceptor. A substrate sits on a susceptor which is at a controlled temperature. The susceptor is made from a material resistant to the metalorganic compounds used; graphite is sometimes used. For growing nitrides and related materials, a special coating on the graphite susceptor is necessary to prevent corrosion by ammonia (NH₃) gas

• Gas inlet and switching system. Gas is introduced via devices known as 'bubblers'. In a bubbler a carrier gas (usually nitrogen or hydrogen) is bubbled through the metalorganic liquid, which picks up some metalorganic vapour and transports it to the reactor. The amount of metalorganic vapour transported depends on the rate of carrier gas flow and the bubbler temperature, and is usually controlled automatically and most accurately by using a Piezocon type vapour control system. Allowance must be made for saturated vapours.

• Pressure maintenance system

• Gas Exhaust and cleaning System. Toxic waste products must be converted to liquid or solid wastes for recycling (preferably) or disposal. Ideally processes will be designed to minimize the production of waste products.
• Fuente: http://en.wikipedia.org/wiki/Metalorganic_vapour_phase_epitaxy
• Nombre: Franklin J. Quintero
• Asignatura:EES

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Hydride Vapour Phase Epitaxy (HVPE)

Oxford Instruments recently acquired Technologies and Devices Inc (TDI). Based at Silver Spring, Maryland, USA, TDI are a world leading company in the development of Hydride Vapour Phase Epitaxy (HVPE) processes and techniques for the production of novel compound semiconductors such as GaN, AlN, AlGaN, InN, InGaN. These materials are used in a variety of applications, the primary ones being solid state lighting, short wavelength optoelectronics and RF power electronics.
Nitride-based Templates by TDI
The company produces a wide range of materials on different substrates, including the following:

Templates	Sizes	Applications
GaN on Sapphire	2" to 4"	Blue and White LED applications
AlN on Silicon Carbide	2" to 4"	Typically used for RF electronic devices such as HEMT
AlGaN on Sapphire	2" or 3"	Optoelectronic devices operating in UV spectral region
InN on Sapphire	Research grade available in 2"	For work on sensors and high frequency electronic devices
InGaN on Sapphire	2" for Green LED	Green LED and green laser Developments

The HVPE Process
In the HVPE process, Group III nitrides (e.g., GaN, AlN) are formed by reacting hot gaseous metal chlorides (e.g., GaCl or AlCl) with ammonia gas (NH3) (Refer to diagram below). The metal chlorides are generated by passing hot HCl gas over the hot Group III metals. All reactions are done in a temperature controlled quartz furnace.
e.g., Hot HCl (g) + Ga (l) ------> GaCl (g)
GaCl (g) + NH3 (g) -------> GaN (s) + HCl (g) + H2 (g)
The GaN or AlN templates have been grown on substrates such as SiC or sapphire. p-type GaN or AlN can be achieved by using Mg during the process and n-type by silane gas with Argon as the carrier gas.
Advantages of HVPE
Developed in the 1960s, it was the first epitaxial method used for the fabrication of single GaN crystals. One of the key features of the technique is its high growth rate (at up to 100 µm per hour) which is almost two orders of magnitude faster than typical MOCVD and MBE processes.
The technique is able to produce crack-free, high quality GaN epitaxial layers (e.g., a typical dislocation density can be as low as 107/cm3 for a 10 µm thick GaN template on sapphire.) Figure 1 shows the X-ray diffraction of a 10 µm thick GaN template on sapphire. The narrow FWHM of 250 arcsec measured at w-scan (0002) peak demonstrates excellent material quality.
Another advantage of HVPE is its ability to grow thick, high quality of AlGaN and AlN for use in optoelectronic and RF electronic devices. The technique has been demonstrated by TDI to grow thicker high quality AlGaN-based active regions of shorter wavelength emitters, which have a high radiative recombination efficiency – an essential feature for high-efficiency UV LEDs. Unlike MOCVD, the HVPE process does not involve metalorganics, thus providing a 'carbon-free' environment for epitaxial growth. In addition, the use of gaseous hydrogen chloride also provides an impurity 'self-cleaning' effect, which results in epitaxial layers with low background impurities and more efficient doping level.
TDI has demonstrated the industry's first HVPE-grown, multilayer, submicron AlGaN/GaN heterostructures.

Fuente: http://www.oxford-instruments.com/products/etching-deposition-growth/processes-techniques/hvpe/Pages/hydride-vapour-phase-epitaxy.aspx
Nombre: Franklin J. Quintero
Asignatura: EES

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Repeated temperature modulation epitaxy for p-type doping and light-emitting diode based on ZnO

Since the successful demonstration of a blue light-emitting

diode (LED)1, potential materials for making short-wavelength

LEDs and diode lasers have been attracting increasing interest

as the demands for display, illumination and information storage

grow2–4. Zinc oxide has substantial advantages including large

exciton binding energy, as demonstrated by effi cient excitonic lasing

on optical excitation5,6. Several groups have postulated the use of

p-type ZnO doped with nitrogen, arsenic or phosphorus7–10, and

even p–n junctions11–13. However, the choice of dopant and growth

technique remains controversial and the reliability of p-type ZnO is

still under debate14. If ZnO is ever to produce long-lasting and robust

devices, the quality of epitaxial layers has to be improved as has been

the protocol in other compound semiconductors15. Here we report

high-quality undoped fi lms with electron mobility exceeding that

in the bulk. We have used a new technique to fabricate p-type ZnO

reproducibly. Violet electroluminescence from homostructural

p–i–n junctions is demonstrated at room-temperature.

Extending semiconductor devices to new compounds

has produced great benefi ts to human life, as exemplifi ed by

modern optoelectronic and high-speed electronic devices for

communications attained by GaAs-based III–V compounds16,

and blue LEDs and lasers realized by GaN (ref. 1). Even widerbandgap

compounds such as diamond, AlxGa1–xN, BN (refs 2–4)

and ZnO have attracted considerable interest for applications in

ultraviolet LEDs and lasers and electronic devices durable at highpower

and/or high-temperature operation. Among them, ZnO has

the following advantages for LEDs and lasers. First, the exciton

binding energy in ZnO is as large as 60 meV and can be increased

to over 100 meV in superlattices17. This exciton stability provides

opportunities for making highly effi cient lasers operable at roomtemperature5,6.

Second, it is possible to tune the bandgap from

3 eV to 4.5 eV in MgxZn1–xO and Zn1–xCdxO alloy fi lms with quite

small lattice mismatch between the two different compositions18–20.

This advantage makes it possible to realize strain-free and highquality

quantum wells. Third, large and high-quality single-crystal

wafers are commercially available.

To harvest these advantages in real devices, a reliable technique

for fabricating p-type doping needs to be established. Compared with

other II–VI semiconductors and GaN, it has been diffi cult to dope

ZnO to produce a p-type semiconductor because of a strong selfcompensation

effect arising from the presence of native defects or

hydrogen impurities21,22. The fi rst p-type ZnO was claimed in fi lms

made by vapour-phase transport7 in NH3, followed by molecularbeam

epitaxy (MBE) using an atomic nitrogen source9. These fi lms

had a hole concentration of 1016–1017 cm–3. There are other claims8,10

that As or P doping can achieve hole concentrations higher than

1 × 1018 cm–3. In the former case, it can be naturally understood that

nitrogen replaces oxygen to generate holes because of similar ionic

radius. But the latter cases have raised the following questions14.

The ionic radii of P and As seem to be too large to occupy the oxygen

site within the wurtzite host lattice and thus to serve as acceptors.

Also, from a comparison with established properties of p-type GaN

(ref. 12), such a high concentration of holes would not be expected

in ZnO. None of the studies mentioned addressed signifi cant effort

to growing the high-quality undoped fi lms that generally serve as

a starting point for reliable doping in semiconductors. Here we

propose a repeated temperature modulation (RTM) technique as a

reliable and reproducible way to fabricate p-type ZnO.

Thin fi lms of ZnO and junction devices were grown by laser

MBE using nitrogen as a p-type dopant (see Methods). Figure 1a

shows a typical intensity oscillation of refl ection high-energy

electron diffraction (RHEED) observed during undoped ZnO fi lm

growth at a temperature (Tg) of 950 °C. This layer-by-layer growth

mode has become possible with our development of an atomically

smooth ZnO buffer layer on ScAlMgO4 (SCAM) substrate produced

by high-temperature annealing23,24. The undoped ZnO fi lms

show excellent optical and electronic properties. Figure 1b shows

a photoluminescence spectrum (inset) and temporal variation

of the photoluminescence intensity taken at room-temperature.

The lifetime (τPL) of free-exciton emission reaches 2.5 ns. This τPL is

much longer than the value in ZnO crystal (1 ns)25 or high-quality

GaN single crystal (0.86 ns)26. Such a long lifetime indicates very

low density of non-radiative defects and negligible carrier trapping

to deep radiative defects, such as the green luminescence band

frequently seen in poor-crystallinity crystals and fi lms. Figure 1c

shows the temperature dependence of residual electron density

(n) and Hall mobility (μ) for an undoped ZnO fi lm. The value of

n is about 1 × 1016 cm–3 at room temperature and decreases with

decreasing temperature, with an activation energy of about 60 meV.

The value of μ is 300 cm2 V–1 s–1 and 5,000 cm2 V–1 s–1 at 300 K and

100 K, respectively, surpassing the best value for a ZnO bulk single

crystal27. Therefore, we conclude that ZnO fi lms grown at this high

temperature, Tg, on atomically smooth buffer layers can serve as an

arena for testing acceptor doping.

However, nitrogen, as one of the most promising acceptor

impurities, cannot be incorporated into ZnO at such a high Tg.

Nitrogen concentration (CN) decreases from a few times 1020 cm–3 at

Tg = 450 °C to a few times 1018 cm–3 at Tg = 700 °C, and at Tg = 950 °C

it is lower than the detection limit (mid-1017 cm–3)28. To solve this

dilemma, we have developed RTM to satisfy both high crystallinity

and high CN. We repeated a growth sequence in which a nitrogendoped

ZnO (ZnO:N) layer 10–15 nm thick is deposited at low

temperature (TL), followed by rapid ramp to high temperature (TH),

and growth of a 1-nm-thick layer at TH. Figure 2a shows an example

of the time variations of RTM. During the processes, the RHEED

pattern switched between two states as shown in the insets of Fig. 2b:

streaks (left) at TL and spots (right) at TH. The streak length, defi ned

by the full width at half maximum of streak intensity peak in the

vertical direction (shown on the insets), is plotted as a function of

time in Fig. 2b. During deposition at TL, the surface became gradually

rougher because of the low surface diffusivity of precursors.

On rapid ramping to TH, the surface smoothness recovered quickly.

Part of the time variation in Fig. 2b is magnifi ed in Fig. 2c (red)

together with the RHEED intensity oscillation (black). The two

oscillation patterns have opposite phase, indicating that initial TL

layers grown on atomically fl at TH layers grew in layer-by-layer mode

even at such a low temperature. Figure 2d shows an atomic force

microscopy image for a ZnO:N fi lm. The surface is composed of

atomically fl at, wide terraces and 0.26-nm-high islands. By using RTM, we

can grow ZnO:N fi lms with CN ranging from several times 1020 cm–3

to a few times 1018 cm–3 by tuning TL from 400 °C to 600 °C.

Among these fi lms, those grown at TL = 400 °C and TH = 950 °C

reproducibly showed p-type conduction. It is worth mentioning that

TH was chosen to satisfy the condition that hydrogen is completely

extracted from the fi lm29. Because residual hydrogen in the growth

chamber is incorporated in the fi lms during growth21, hightemperature

annealing may be essential not only for annihilating

non-equilibrium defects but also for removing hydrogen, if any, to

activate the acceptors.

The ZnO:N fi lms prepared by RTM have the same in-plane lattice

constant as that of the buffer layer. The out-of-plane lattice constant

is slightly expanded (+0.02%) compared with the bulk value.

The rocking curve width is as narrow as that of the undoped ZnO

fi lms, indicating that the structural quality of the p-type ZnO is very

high. The inset of Fig. 3 shows a set of raw data for the Hall resistance,

indicating p-type conduction as evidenced by the positive slope.

Figure 3 shows hole concentration as a function of temperature for

a p-type ZnO:N fi lm with CN = 2 × 1020 cm–3. From these data, we

can deduce an activation energy EA of 100 meV and compensation

ratio ND/NA ≈ 0.8. Here we note that ND/NA obtained in the present

p-type ZnO fi lm is higher than the reported value of 0.1 for p-type

ZnO grown by MBE9. Therefore, there is still room to increase hole

concentration by tuning the RTM growth condition.

The schematic structure of a typical homostructural p–i–n

junction is shown in Fig. 4a, which was grown throughout in layerby-

layer mode keeping an atomically fl at interface. Figure 4b shows

typical current–voltage characteristics. Fairly good rectifi cation

was obtained with a threshold voltage of about 7 V. The threshold

voltage is higher than the bandgap of ZnO (3.3 eV), mainly owing to

the high resistivity of the p-type ZnO layer. The electroluminescence

is measured by feeding in a direct current at room temperature.

This spectrum was measured from the top by detecting the light

escaping from the edge of the top electrode. The electroluminescence

spectrum shows luminescence from violet to green regions with

multi-refl ection interference fringes. We would expect exciton

emission at 3.2 eV from the undoped layer (i-ZnO) as shown in the

inset of Fig. 1b, but the electroluminescence spectrum apparently

shows a redshift. This is partly due to the low hole concentration in

p-type ZnO: electron injection from i-ZnO to p-type ZnO overcomes

hole injection from p-type ZnO to i-ZnO. The photoluminescence

spectrum (black) of a p-type ZnO fi lm is also shown. The higherenergy

side peak around 430 nm in the electroluminescence

spectrum matches well with the photoluminescence spectrum.

Another factor could be that the electroluminescence from active i-

ZnO is partly absorbed in the p-type layer because of a slight redshift

of the absorption edge for this layer. Nevertheless, the signal-tonoise

ratio of the electroluminescence spectrum shown in Fig. 4c is

much better than those reported previously11,30. More importantly,

the RTM technique established for reliable p-type doping of ZnO

enables us to improve the device performance further by optimizing

the growth parameters and device structures.

The next challenge will be to increase the hole concentration

by further optimizing the growth process of p-type ZnO.

Making a p-type (Mg,Zn)O fi lm is also an important challenge, not

only to prevent the majority of electrons from injecting into the

p-type layers, but also to avoid attenuation of band-edge emission

from the i-ZnO.

Figure 1 Thin fi lms of ZnO grown in persisting layer-by-layer mode show

high-quality optical and electronic properties. a, RHEED intensity oscillation

observed during ZnO thin-fi lm growth by laser MBE at a temperature of 950 °C on an

atomically smooth ZnO buffer layer formed on a ScAlMgO4 substrate. The oscillation

period corresponds to a 0.26-nm-thick charge-neutral molecular layer grown by

about 23 laser pulses. The oscillation persisted over fi lm growth to a thickness of

1 μm. The fi lm has the same lattice constants (a = 0.3250 nm, c = 0.5204 nm) as

the bulk values by relaxing the very small lattice mismatch of 0.09% with the SCAM

substrate in the high-temperature annealed ZnO buffer layer. The full-width at halfmaximum

of the rocking curve at the (002) refl ection is less than 18 arcsec, which

is close to the instrumental resolution. The growth direction is identifi ed, through

experiments similar to those reported previously32, to be [0001–] of the wurtzite

structure (oxygen face). b, Photoluminescence (PL) spectrum (inset) and temporal

decay of the luminescence intensity at the peak energy (indicated by triangle)

for an intrinsic ZnO fi lm. The dashed peak represents the temporal evolution of

the excitation laser pulse (Ti: sapphire, 242 nm, 30 mW, 80 fs). The lifetime of

free exciton emission exceeds 2.5 ns, indicating the high quality of the sample.

c, Temperature dependence of electron mobility and carrier concentration for an

undoped ZnO fi lm (solid circles) and a ZnO bulk single-crystal (open circles, after

Look et al.27). To ensure that the measurements extract the intrinsic properties of

the ZnO fi lm, a fairly thick fi lm (1 μm) was grown on a semi-insulating Mg0.15Zn0.85O

buffer layer which was annealed to prepare an atomically smooth surface before

the ZnO deposition23.

Figure 2 Atomically smooth ZnO fi lms doped with nitrogen can be grown by

a repeated temperature modulation technique. a, Temporal variation of growth

temperature, switching between TH and TL, during ZnO thin-fi lm growth. Layers

of ZnO:N with high nitrogen concentration (CN) were deposited at TL in the period

coloured blue. The layers were annealed and additional ZnO:N layers with low CN

were grown in the period coloured red in order to activate nitrogen as an acceptor

and recover surface smoothness, respectively. b, Typical RHEED patterns observed

during TL (left) and TH (right) periods are shown in the insets, representing rather

rough and atomically smooth surfaces by streaky and spotty patterns, respectively.

Temporal variation of streak length (defi ned by broken lines) is plotted as a measure

of the surface roughness. c, RHEED intensity (black) and streak length (red) are

plotted for the initial growth of high CN layer at TL as denoted by circle in b. Clear

oscillations having a half-phase shift confi rm the layer-by-layer growth mode.

d, An AFM image of the surface for a 500-nm-thick ZnO:N fi lm. The step height

corresponds to the thickness of a charge-neutral molecular layer of ZnO (0.26 nm).

Figure 3 Temperature dependence of hole concentration (p) in a p-type

ZnO doped with nitrogen. (Nitrogen concentration CN = 2 × 1020 cm–3.) A typical

variation of the Hall voltage during a magnetic fi eld scan is shown in the inset.

The hole mobility varies from 5 cm2 V–1 s–1 at 350 K to 8 cm2 V–1 s–1 at 300 K.

Activation energy EA and compensation ratio ND/NA are deduced to be 100 meV and

0.8, respectively, from the linear fi tting9 of p = (ND/NA – 1)(gA1/gA0)NvT3/2 exp(–EA/kBT),

where gA0 = 4 and gA1 = 1 are the unoccupied and occupied state degeneracies,

respectively, NvT3/2 = 2(2πmhkBT/h2)3/2 is the density of states in the valence band,

where kB is Boltzmann's constant and h denotes Planck's constant, and mh is

assumed to be 0.9m0 with m0 being free electron mass.

Figure 4 Zinc oxide homostructural p–i–n junction shows rectifying

current–voltage characteristics and electroluminescence (EL) in forward

bias at room-temperature. a, The structure of a typical p–i–n junction LED.

b, Current–voltage characteristics of a p–i–n junction. The inset has logarithmic scale

in current with F and R denoting forward and reverse bias conditions, respectively.

c, Electroluminescence spectrum from the p–i–n junction (blue) and

photoluminescence (PL) spectrum of a p-type ZnO fi lm measured at 300 K.

The p–i–n junction was operated by feeding in a direct current of 20 mA.

fuente: http://lib.semi.ac.cn:8080/tsh/dzzy/wsqk/Nature%20Material/nmat4-42.pdf

Nombre: Franklin J. Quintero

Asignatura: EES

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