Crecimiento Solvotermal de ZnO

Progress in Crystal Growth and Characterization of Materials

52 (2006) 280e335www.elsevier.com/locate/pcrysgrow

Solvothermal growth of ZnO

Dirk Ehrentraut a,*, Hideto Sato b, Yuji Kagamitani a, Hiroki Sato a,Akira Yoshikawa a, Tsuguo Fukuda a

a Institute of Multidisciplinary Research for Advanced Materials, Tohoku University, 2-1-1 Katahira,

Aoba-ku, Sendai 980-8577, Japanb Murata Mfg. Co., Ltd., 1-10-1 Higashikotari, Nagaokakyo, Kyoto 617-8555, Japan

Abstract

The growth of ZnO single crystals and crystalline films by solvothermal techniques is reviewed. Larg-est ZnO crystals of 3 inch in diameter are grown by a high-pressure medium-temperature hydrothermalprocess employing alkaline-metal mineralizer for solubility enhancement. Structural, thermal, opticaland electrical properties, impurities and annealing effects as well as machining are discussed. Poly-and single-crystalline ZnO films are fabricated from aqueous and non-aqueous solutions on a variety ofsubstrates like glass, (100) silicon, a-Al2O3, Mg2AlO4, ScAlMgO4, ZnO and even some plastics at tem-peratures as low as 50 �C and ambient air conditions. Film thickness from a few nanometers up to sometens of micrometers is achieved. Lateral epitaxial overgrowth of thick ZnO films on Mg2AlO4 from aque-ous solution at 90 �C was recently developed. The best crystallinity with a full-width half-maximum from the (0002) reflection of 26 arcsec has been obtained by liquid phase epitaxy employingalkaline-metal chlorides as solvent. Doping behavior (Cu, Ga, In, Ge) and the formation of solid solutionswith MgO and CdO are reported. Photoluminescence and radioluminescence are discussed.� 2006 Elsevier Ltd. All rights reserved.

PACS: 68.35.Dv; 68.55.Jk; 68.55.Nq; 71.55.Gs; 78.55.Et; 81.05.Dz; 81.15.Lm; 81.10.Dn; 81.20.Wk

Keywords: A1. Solvents; A1. Substrates; A2. Hydrothermal crystal growth; A3. Liquid phase epitaxy; B1. Zinc com-

pounds; B2. Semiconducting IIeVI materials

* Correponding author. Fax: þ81 22 217 5102.

E-mail address: [email protected] (D. Ehrentraut).

0960-8974/$ - see front matter � 2006 Elsevier Ltd. All rights reserved.

doi:10.1016/j.pcrysgrow.2006.09.002

mailto:[email protected]

http://www.elsevier.com/locate/pcrysgrow

281D. Ehrentraut et al. / Progress in Crystal Growth and Characterization of Materials

52 (2006) 280e335

1. Introduction

Over the last few years an enormous effort has been dedicated towards the film growth ofZnO in order to meet the needs of a large variety of applications [1,2]. The quality of ZnO filmshas significantly improved and recently a room temperature (RT) ZnO-based light-emitting di-ode (LED) has been demonstrated by Tsukazaki et al. [3]. Often due to the lack of high-qualityZnO substrates, most of the device structures are based on heteroepitaxial ZnO films, i.e. ZnOdeposited on foreign substrates with different crystallographic and thermal properties. The re-sulting disadvantages in ZnO films on a-Al2O3 or GaN involve, i.e. a temperature-dependentlattice misfit that often makes buffer layers indispensable, and the out-diffusion of ions likeAl and Ga from the substrate into the ZnO film [4,5]. This has accelerated the quest forlarge-size ZnO substrates of excellent crystallinity and low defect concentration. The growthhas been carried out by chemical vapor transport (CVT) [6], pressure-melt [7], flux (Table 1)and the hydrothermal technique (Table 2). The largest high-quality crystals are currently pro-duced by the hydrothermal technique, which now is capable of producing specimens 3 inch indiameter (i.e. perpendicular to the h0001i direction).

Nearly lattice matched solid-solution films can be grown within a large area of the wurtzite-type MgOeZnOeCdO system [8], e.g. the in-plane lattice mismatch between (0001)Mg0.2Zn0.8O and (0001) ZnO is as small as 0.13% [9]. The growth of ZnO layers is dominatedby technologies employing the vapor phase, i.e. pulsed laser deposition (PLD), molecular beamepitaxy (MBE), Metal-organic chemical vapor deposition (MOCVD), and physical vapor depo-sition (PVD). A very comprehensive overview is given by Triboulet and Perriere [2]. Vapor-phase processing and properties of ZnO-based films has recently been reviewed by Peartonet al. [1]. Ohtomo and Tsukazaki [10] report on state-of-the-art in growth of thin films andsuperlattices based on ZnO by PLD.

However, due to advantages like simple equipment, low temperature and ambient pressure,vicinity to the thermodynamic equilibrium, etc., growth from the liquid phase is attractive andhas recently been demonstrated for different fluorides and oxides including ZnO [11e19]. De-spite all efforts, fabrication of doped ZnO often remains a challenge where solvothermal (i.e.using a solvent at elevated temperature at which the solvent is in its liquid phase state) tech-niques may possess an advantage over vapor-phase techniques. By growth from a liquid phasethe solute species are transported as metastable molecules and subsequently decomposed ata growing surface due to thermodynamic instability. The doping species is required to be incor-porated in the correct valence state on either the zinc or the oxygen lattice site. In vapor growthtechnologies, additional activation of the doping species is therefore necessary. On the otherhand, higher doping levels can be achieved by some vapor-phase growth techniques due to ther-modynamically off-equilibrium deposition conditions.

This review reports on the fabrication of ZnO crystals from non-aqueous solvents and, ingreater detail, by the high-pressure hydrothermal technique in Section 3. However, it is regret-table that much of the initial work done in former Soviet Union laboratories is not available.

The characteristics of ZnO wafers machined from hydrothermal ZnO are analyzed inSection 3.

Section 4 comprises the solvothermal film growth from low-temperature aqueous solutionsand by water-free liquid phase epitaxy (LPE) employing chloride solvents. The LPE films arecharacterized in detail.

The constantly increasing field of ZnO-based nanocrystals [20] and nanostructures [21] willnot be touched by this review although growth from solution is a major technology there as well.

Ta

Zn

G Results Refs.

So Crystal size, growth rate, remarks

22 1e5 cm Long in h11e20i direction; habit

change of crystal: growth at below 1050 �C

yields drum-shaped crystals, platelets for the

growth at 1150 �C; not possible to chemically

separate ZnO from the flux.

[23]

Zn [24]

20 5� 5� 3 mm3; 0.5 K cm�1�DT� 3 K cm�1;

Inclusions parallel (0001) planes; high DT of

3 K cm�1 yields nucleation only on the bottom

of the crucible; air jet was used; strong

evaporation of PbF2 of up to 40 wt%.

[25]

11 Pale green, clear crystals about 8 mm across;

polycrystalline.

[26]

42 Platelets, 20 mm across.

Zn

76 TSSGb; polycrystalline boule 22 mm diameter�4 mm; single crystals up to 10� 5� 2 mm3;

0.8% V and 0.7 % Mo contamination

[27]

52 TSFZc; rod size 4� 12; grains up to 2� 2� 1 mm3;

1.7% V contamination.

b

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ble 1

O crystals grown from non-aqueous solutions

rowth conditions

lute Solvent T (�C);

DT (K)

Cooling rate

(K h�1)

Atm. Seed Crucible

g ZnO 200 g PbF2 800e1150 1e10 Air, O2 None Pt, 100 ml

O PbF2 1000 5.4

mol% ZnO 80 mol% PbF2 800e1250 2e4 Air None Pt, 100 &

250 ml

7.3 g ZnO 92.4 g V2O5 900e1300 1.2 Air None Pt

.9 g ZnO 37.3 g Zn3P2O8�4H2Oþ 42.9 g V2O5

980e1330 1

O:

e80 mol% 20e24 mol%

V2O5þB2O3

1150 2e5 Air Pt wire Pt, 55�40 mm2

mol% 48 mol% MoO3þV2O5

Molten zone 0.5e1 mm h�1

growth rate

None None

a n.r. e Not reported.

TSSG e top-seeded solution growth.c TSFZ e traveling-solvent floating zone.


52 (2006) 280e335

2. General aspects of growth from solution

In general, a solution consists of the solute and a solvent, which serves to transport thedissolved solute to the growth interface of the crystal. The chemical nature of the solvent,mainly determined by its bonding type, is a key point for chemical interaction with the sol-ute. In some cases like for the growth of ZnO from aqueous solutions (Sections 3.2.1 and4.1), enhancement of the low solubility by applying mineralizers is required. The ideal prop-erties of a solvent to be applied in high-temperature solution growth have been summarizedby Elwell and Scheel [22] and may as well apply to processes at low temperature: (a) highsolubility for the solute, (b) appreciable change of solubility with temperature, (c) rather lowmelting point, (d) low vapor pressure to avoid changes in the solventesolute ratio, (e) therequired crystal phase should be the only stable solid phase, (f) low dynamic viscosity inthe range between 1 and 10 mPa s, (g) low reactivity with the material of the growth vessel,(h) appropriate physical density (i) the ease of separation of the grown crystal from the sol-vent by chemical or physical means, (j) availability in high purity at reasonable cost, and (k)low toxicity.

General knowledge about solution growth is the fact that highest crystallinity can beachieved because the growth is conducted in the vicinity of the thermodynamic equilibrium.This is the point where neither growth nor dissolution of a given crystal occurs. If we increasethe concentration of the solute relative to the concentration needed to saturate the solvent, i.e. toestablish thermodynamic equilibrium, a supersaturation is built up. Finally, deposition of thesolute on a seed crystal (¼crystal growth) is achieved upon reaching the threshold for supersat-uration. Control of the supersaturation is usually made by managing the concentration of thesolute species through the absolute temperature of the solution and mass transport. Controlof the geometrical direction of the supersaturation to trigger crystal growth on the seed crystalcan best be achieved by establishing a temperature gradient between seed crystal and its sur-rounding solution in the reactor. This is basically an engineering problem of designing the suit-able reactor geometry.

The ecological aspect of the discussed solvothermal syntheses is that this technology is en-vironmentally benign. Solvents like water or alkaline-metal chlorides are easy to recycle.Growth temperatures are rather low and often air atmosphere is employed, which also contrib-utes to keep production costs low as well as the ease of maintenance of the rather simple

Table 2

An overview over some conditions and results for the growth of SiO2 and ZnO from the presently largest autoclaves in

production. Calculation of crystal weight, yield, and growth speed relates to a growth period of 100 days [34]

Parameter SiO2 ZnO

Autoclave I.D. 0.65 m �0.2 ma

Autoclave I.L. 14 m �3 ma

Volume 4.6 m3 �0.2 m3

Seed size 70� 45� 230 mm3 z50 mm diametera

Seeds per batch 1400 112a

Weight of crystal 1700 g 320 g; 20 mm thick

Total yield per batch 2300 kg 36 kg

Growth speed (c-axis) 500e600 mm/day 200 mm/day

3-runs-per-year yield 6900 kg 108 kg

a Value may marginally differ.

284 D. Ehrentraut et al. / Progress in Crystal Growth and Characterization of Materials52 (2006) 280e335

equipment. A high throughput is made possible by using large growth vessels like those usedfor the production of hydrothermal quartz (Fig. 2).

3. Growth of ZnO crystal

3.1. Non-aqueous solutions

Attempts have been made to grow ZnO crystals from the following molten salts: PbF2,Zn3P2O8 � 4H2O, V2O5, MoO3, B2O3, and mixtures thereof [23e27]. Table 1 comprisesgrowth conditions and results. The growth temperatures are �800 �C. The crystals grownfrom PbF2 are impossible to separate from the solidified flux [23]. Often the growth of plateletsis reported [23,26], which is likely due to impurity effects caused by the solvent. This modifi-cation of the growth mechanism on some defined crystallographic directions by impurities isknown from a variety of mainly solution-grown crystals. It was explained by a substantialchange of the average binding forces operating along the surface between particles of the sur-face layer due to adsorbed impurities [28]. The pale green color of the crystals in Ref. [26] islikely due to impurities derived from the solvent. Also, inclusions parallel to (0001) planes havebeen reported [25]. Temperatures higher than 1300 �C led to strong evaporation of PbF2. Theuse of a mixture of Zn3P2O8� 4H2O and V2O5 by Wanklyn [26] gained 20 mm large platelets.It was reported that the mixtures did not adhere to the crystals. Separation of the liquid solutionfrom the grown crystals was carried out by pouring it off the crystals. All crystals were single-phase ZnO. The resistivity was measured as 0.3 U cm. This is comparable to values obtainedfrom pressure-melt grown ZnO, compare Fig. 19c, and hints to higher impurity levels.

The top-seeded solution growth (TSSG) and traveling-solvent floating zone (TSFZ) tech-niques were applied to grow larger crystals from mixtures of ZnO with V2O5þ B2O3 andMoO3þV2O5 [27]. TSSG comprises crystal growth on a seed immersed in a solution. In orderto control the temperature field around the growing crystal, the seed crystal is simultaneouslyrotated and pulled from the solution at approximately at the rate at which the crystal is growing.In TSSG of ZnO a platinum wire served for seeding and polycrystalline aggregates with single-crystalline regions of 10� 5� 2 mm3 were obtained. The crystal pulling and solution coolingrates were 0.5e1 mm h�1 and 2e5 �C, respectively. Crystal rotation rate was 20 rpm. Growthby TSFZ involves providing a feeding rod and a seed rod concentrically arranged to give a pointof contact between both rods. At this contact point, the feeding rod is melted by an externalheater. Successively, this molten zone is caused to travel toward the end of the feeding rodby moving the heater. After a first melting at a travel speed of the solution zone of15e20 mm h�1, the growth was carried out at rates of 0.5e1 mm h�1. Feedstock and growingcrystal were rotated at about 20 rpm in opposite directions. Grains up to 2� 2� 1 mm3 wereobtained. In either case, however, a large amount of metal impurities up to 1.7% V wasmeasured and contamination from the solvent was considered to be a serious issue [27].

Our own experiences with non-aqueous solution systems comprise the system PbOeZnOand employment of Zn3(PO4)2eLi2O. Whereas the first case showed the above-mentionedproblem of ZnO separation from the PbO solvent after the growth and was therefore excludedfrom further investigation, the latter system was tried for a while. Following reaction path wasused:

Zn3ðPO4Þ2 þ 3Li2O/3ZnOþ 2Li3PO4: ð1Þ


52 (2006) 280e335

The formation of phase-pure wurtzite ZnO at temperatures below 1000 �C was confirmed byXRD measurements. However, the whole solution quickly solidified after completion of reac-tion (1). The melting point of Li3PO4 (1205 �C [29]) is considerably higher than of Zn3(PO4)2

(900 �C [29]), which caused the solidification. Attempts to seek a solvent for Li3PO4 weremade among some phosphates like K4P2O7 and KPO3, but the experiments did not lead to sat-isfying results. The path with K4P2O7 would form KZnPO4, which is a competing phase toZnO. The use of KPO3 would result K3PO4þKPO3, which would not solve the problem oflowering the melting point.

In summary, the growth from non-aqueous solutions did not lead to the production of largeZnO crystals of low-impurity concentration and high crystallinity. Reasons why this routefailed can be summarized as: (a) a proper solvent was not available, which would keep the sol-vent-derived impurity level sufficiently low. (b) The separation of grown ZnO from the solventafter the growth process has finished remained an unsatisfactory issue in most cases. (c) Thehydrothermal growth of ZnO was tried at the same time and turned out to be more competitivein terms of crystal size and quality. Supporting is the fact that much knowledge on the hydro-thermal growth technology has been collected during the development of the hydrothermalgrowth of quartz (a-SiO2). This dates back to Spezia in the year 1909 [30] who was first toreport valuable results on seeded growth of quartz under hydrothermal conditions. He alreadyapplied a temperature gradient between the zone for dissolving the feedstock and the growthzone. However, the story of industrial growth of quartz dates back to the 1940s. In 1949, i.e.Buehler and Walker [31] reported the hydrothermal growth of twin-free quartz. In 1953, Walker[32] published the hydrothermal synthesis of quartz crystals weighing over 1 lb (453.6 metricgram), grown at the Bell Telephone Labs. in periods of less than two months. The same paperillustrates two crystals grown over 42 days to a weight of 540 g each. The temperature-difference method and quartz nutrient was employed in autoclaves sizing 10 cm internaldiameter and 122 cm internal length. Walker already pointed out in his paper that crystalsgrown from solutions are likely to be of more perfect quality than those grown from melts.A comprehensive review on the history of quartz crystal growth was recently given by Iwasakiand Iwasaki [33]. Nowadays, the hydrothermal growth of quartz is a routine. In 2004, about1850 tons were produced worldwide with some 700 tons in Japan [34]. Table 2 comparesthe growth of quartz and ZnO from, to our knowledge, the largest autoclaves at present. Theimpressive amount of 2300 kg quartz from 1400 crystals can be produced over a 100-daygrowth run. The same figures for ZnO, 36 kg and about 100 crystals, are still not comparablewith quartz, but really show the potential of the hydrothermal technology in terms of through-put. However, the driving force for scale-up of the hydrothermal growth of ZnO comes from thedemand by industry.

The hydrothermal growth of ZnO and characteristics of the crystal will be treated in greaterdetail in the following section.

3.2. Hydrothermal growth of ZnO

3.2.1. General characteristics of the methodThe term ‘‘hydrothermal’’ is derived from geology [35]. Hydrothermal growth comprises the

use of aqueous solvents and mineralizers under elevated temperature and pressure in order todissolve and recrystallize materials, which are barely soluble under ordinary conditions.Mineralizers are particularly important since they serve to establish a suitable solubility ofthe solute because most of the species are rather poorly soluable in water.


A large variety of crystals have been grown with a-quartz (SiO2) being the most prominentone and most of the technological developments in the hydrothermal technique are closely re-lated with the development of the a-quartz technology. A detailed description on the hydrother-mal technology with examples for many crystals is given by K. Byrappa [35].

Features of the hydrothermal method comprise: (a) use of a closed high-pressure growth ves-sel (autoclave); (b) use of a solvent; (c) use of solubility increasing mineralizers; (d) employ-ment of a precursor; (e) employment of seed crystals; (f) a temperature gradient between theprecursor-containing dissolution zone and the growth zone with the seed crystals;(g) DT z 0 at the interface between the growing crystal and the solution, which is why theconcentration of structural crystal defects is smaller than for melt-grown crystals; (h) saturationof the solute while the seed crystal is already in contact with the undersaturated solution.

The hydrothermal growth of ZnO requires the use of water in its supercritical state, which isachieved at its critical point with a critical temperature and pressure, Tc¼ 374 K andpc¼ 22.1 MPa, respectively. Fig. 1 shows the simplified peT diagram for water at constant vol-ume. The existence region of supercritical water (SCW) covers the upper right hand side abovethe critical point. The peT region at which large hydrothermal ZnO crystals are grown is rep-resented by the rectangle, which stretches between the pressure of 70e255 MPa at temperaturesof 300e430 �C, see also Table 3.

SCW is characterized by an enhanced acidity, reduced density (0.05e0.2 g cm�3) and lowerpolarity in comparison to water under normal pressure and temperature. In fact, SCW is an al-most non-polar fluid due to reduced dielectric strength to values of 1e3 (water under ambientconditions shows 78) [36]. The diffusivity of SCW is strongly increased as well as the misci-bility with gases. SCW exhibits reduced molecular ordering with less effective hydrogen bond-ing [37]. The enhanced acidity favors ionic processes, such as the dissolution of ZnO. However,the solubility of ZnO in SCW remains insufficient and makes use of mineralizers necessary.

Mineralizers serve to increase the solubility of ZnO in SCW by forming metastable com-pounds between them and ZnO, which later decompose at a growing crystal face to deliverZn2þ and O2� which are incorporated as the ZnO lattice. Typical mineralizers are LiOH,NaOH, KOH, Li2CO3, and H2O2, see also Table 3. The best solvents for ZnO are a mixture

Fig. 1. Schematic pressureetemperature diagram for water at the condition of constant volume.

Table 3

Growth conditions of hydrothermal ZnO crystals

Result Refs.

Baffle Crystal size, growth rate

n.r. Prior growth seeds were etched

in HCl and NaOH solution.

[38]

Pt 10 mm B after 14 days [40,41]

n.r. n.r. [42]

n.r. (0001): 0.25e0.38 mm/day;

ð1011Þ: 0.12e0.25 mm/day

[43]

5% (0001): 0.35 mm/day

(10-day growth run) and

0.3 mm/day (30-day growth run)

[44]

5% Average {0001} z 0.25 mm/day [45]

0.53e0.75 mm/day

n.r. 15� 15� 8 mm3; (0001):

0.45 mm/day; ð0001Þ:0.22 mm/day; Perpendicular

to c: 0.35 mm/day

[46]

Ag,

5%

(0001) and ð1011Þ:z0.2 mm/day

[47]

n.r. (0001): averaged 0.25 mm/day

for 30 days growth run

[48]

n.r. Nitrogen doping

up to 8� 1018 atom cm�3

confirmed by inert gas fusion

analysis

[49]

(continued on next page)

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52(2006)

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Growth conditions Autoclave

Precursor Mineralizer T (�C);

DT (K)

P

(MPa)

Seed Filling

(%)

Crucible

lining

ZnO KOH, LiOH, NaOH,

NH4OH

230e300; 5e75 50e350 different

orientations

n.r.a n.r.

Sintered ZnO,

1 B� 3 mm33 M KOHþ 1 M LiOHþ0.1 M H2O2

370e415; 10 70e100 (0001) n.r. Pt

n.r. 3 N NaOHþ 1 N KOHþ0.5 N Li2CO3

345e355; 10 n.r. (0001)

20 mm2�0.5 mm

80 Pt

n.r. 1e2 M NaOH 387e430; 8e30 145e255 (0001),

ð1011Þ70e85 n.r.

ZnO (grain size

larger U.S.

Standard

Sieve Size

No. 10)

5.1 N, i.e. 5.45 molal

KOHþ 0.7 molal LiOH

340e385 (best

at 353); 14

(0001) 82e87;

best at 83

Ag

Hydrothermal

recrystallized ZnO

5.1 M KOH 340e350; 10e15 �55 (0001) 83e85 Ag

5.1 M KOHþ 2 M LiOH 386e428; 25e30 116e169 79e86.1

Pressed and

sintered

ZnO pellets

6 M KOHþ 1 M LiOH Nutrient 365; 10 n.r. (0001)

7 B mm283 Pt

Pressed and

sintered pellets,

1 B� 3 mm3

3 M KOHþ 1e2 M LiOH 360e380;

z10e25

z100 (0001) n.r. Ag

ID 35 mm

L 350 mm

Sintered powder

99.99% ZnO

4 M KOHþ 4 M NaOHþLi2CO3

Nutrient 355; 10 n.r. (0001) 80 Pt

Sintered powder

99.99%, particle

size< 3 mm

3 M NaOHþ 1 M KOHþ1 M LiOH;

350e365; 15 20 (0001) 75 Pt

20 vol% N H4OHþ80 vol% 5 M KOH

475e490; 15 136 70

Table 3 (continued )

Growth conditions Autoclave Result Refs.

cible

g

Baffle Crystal size, growth rate

D

m,

16e

m3

n.r. Up to 5.5 mol% Mg in the

ZnO crystal; ZnO film formed

on the surface of the crystals

during cooling to room

temperature.

[50]

ml

7e15% (0001): max. 2.053 mm/day;

ð0001Þ: max. 1.097 mm/day;

Both for Li-free solvent

(5.15 M KOH)

[51]

Pt 50� 50� 15 mm3 [52,53]

50 mm,

00 mm

5e6% n.r. [54]

on,

ml

me

n.r. 16 h growth

time, precipitates

[55]

on Not

used

Needles, 1 mm in (0001)

from EDTA; duration 7e11

days

[39]

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Precursor Mineralizer T (�C);

DT (K)

P(MPa)

Seed Filling

(%)

Cru

linin

Sintered powder 3e5 M KOHþ�20 wt%

MgO (99.999%)

675e650; �25 102.5e

123.7

Single

crystals

used

45e60 Pt I

14 m

vol

33 c

Fine crystalline

ZnO powder

5e5.4 M KOHþ0.8e1.2 M LiOH

250e300; 20e80 15e140 (0001),

ð1010Þ,ð1011Þ

n.r. Ag,

500

Sintered ZnO

poly-crystals

3 M KOHþ 1 M LiOH 300e400 80e100 (0001) n.r. Pt

Sintered ZnO,

99.99%, granules,

1e6 mm

LiOH, NaOH, KOH 300e400; 10e15 100e

150

VP- and

hydrothermal

n.r. ID

L 6

Zn(OAc)2� 2H2Ob,

Co(OAc)2� 4H2Oc80 ml KOH 240 n.r. n.r. n.r. Tefl

100

volu

Zn(OAc)2� 2H2O DTPAd, EDTAe, TEPfþKOH

200 n.r. None n.r. Tefl

a n.r. e not reported.b Zn(OAc)2� 2H2O e zinc acetate.c Co(OAc)2� 4H2O e cobalt (II) acetate.d DTPA e diethylenetriaminepentaacetic acid.e EDTA e ethylenediaminetetraacetic acid.f TEP e tetraethylenepentamine.


of KOH and LiOH. Pure KOH gives rise to a high growth rate which leads to low crystal qualityand the growth process is hard to control. By contrast, LiOH alone is too weak to significantlyincrease the solubility of ZnO in SCW [38]. Diethylenetriaminepentaacetic acid (DTPA), ethyl-enediaminetetraacetic acid (EDTA) and tetraethylenepentamine (TEP) was employed to growZnO by a low-temperature metal-organic approach [39]. In this case, however, the crystalsize did not exceed 1 mm. Larger ZnO crystals grown in this way have not yet been reported.However, the very low growth temperature is an interesting aspect there.

In Table 3 are listed growth conditions and results obtained by several groups over more thanfour decades or so. The chemical reaction of ZnO and H2O to form ZnO�H2O has already beeninvestigated by Roth and Chall [56] and a phase diagram for the temperature interval 280e320 K was reported as early as 1933 by Huttig and Moldner [57]. However, first growth of largerZnO crystals was published in the 1960s by the group of Kolb and Laudise [43e45,58e60].

Disadvantages of the method are the inability to directly manipulate during a growth run aswell as the lack of visual observation. Temperature changes driven by the heaters assembledaround the autoclave have a long time constant; this prevents sharp temperature changes in thesolution. Therefore, quasi steady-state conditions are applied to hydrothermal crystal production.

3.2.2. Technology for ZnOResearch on the hydrothermal growth of ZnO was carried out mainly by using Morey and

Tuttle type autoclaves [35,44,61,62]. The large capacity autoclaves used for the production ofquartz (Fig. 2) and now for ZnO [63] are basically modified Bridgman autoclaves. A good over-view over the different autoclave types and their specifications is given by Byrappa [35].

Fig. 2. (a) Quartz crystals being withdrawn from the autoclave of 0.65 m inner diameter and 14 m length; (b) crystal

holder with about 25 cm large quartz crystals.

52 (2006) 280e335


Laudise et al. [44] reported on six parameters to be important in the hydrothermal growthof ZnO: (1) a high concentration of alkaline mineralizers is required to achieve a propersolubility of ZnO [64] in SCW while keeping the dynamic viscosity of the solution suffi-ciently low. (2) The temperature difference DT between the dissolution and growth zoneshas to be such as DT� 20 K and control of DT within �3 K is inevitable to suppress thetendency toward spontaneous nucleation and flawed growth in ZnO. This rate-limiting pro-cess was explained in terms of diffusion caused by a high concentration of mineralizers andtheir less effective mixing. (3) Wall nucleation, spontaneous nucleation and flawed growthwere reduced by a two-step warm-up process. Firstly, the long autoclaves were slowlyheated over a period of 24 h. Secondly, further heating was maintained by keeping the con-ditions of DT� 25 K at 150 �C and DT� 20 K above 250 �C. (4) Liþ containing mineral-izers prevented the formation of growth hillocks in ZnO, i.e. free of flaws. (5) Seedsoften possessed a strongly damaged surface as result of mechanical machining, see also Sec-tion 3.2.4. Good-quality ZnO was grown when a surface depth of �50e70 mm was etchedaway from (0001) oriented ZnO seeds. (6) Size of the precursor: very small particle sizeresulted in low growth rates and flawed growth. Best results were found when the precursorwas larger than U.S. Standard Sieve Size No. 10. The use of hydrothermal grown lumpsabout 6 mm in size did not yield an improvement in either growth rate or quality. Theseexperimental findings on the optimum precursor size were recently confirmed by Chenet al. [65] using numerical simulation to analyze the ammonothermal growth of GaN[66], which is very much inspired by the hydrothermal growth of ZnO. The flow of solutionin the precursor stock and the temperature distribution were optimum for a precursor particlesize of 3 mm. In this case, significant convective effects were seen in the precursor stockand the flow was highly three-dimensional. When the particle size was reduced to0.6 mm, a very weak flow was obtained inside the precursor stock and the temperature dis-tribution was controlled by thermal conduction (conduction mode). Dissolution of the pre-cursor and mixing of the precursor-enriched solvent with the solvent above the precursorstock was therefore poor [65].

On the other hand, by increasing the particle size the surface/volume ratio becomes signif-icantly smaller, i.e. the surface available for dissolution by the solvent decreases. Simulta-neously, the temperature distribution is controlled by the solution flow (flow mode).

A basic requirement for the growth from autoclaves is the need for a constant mass flowcirculating from the feedstock to the seed crystal. This is established by proper heater arrange-ments around the autoclave. The inner part contains a baffle, which separates the feedstock(saturation zone; bottom of the autoclave) from the seed crystals (growth zone; upper fractionin the autoclave). The parameters for the baffle are the surface of the bore holes (opening typ-ically around 5e15%) and their geometrical arrangement. If necessary, even several bafflescould be used to get better control of the mass flow. Unfortunately, details on the inner con-struction of autoclaves ready to produce large ZnO crystals cannot be disclosed at present asthey are the key to successful production of high-quality ZnO crystals and therefore intellectualproperties of the ZnO-producing companies. The same is true for the growth process.

The filling of an autoclave describes the amount of liquid solution which is introduced intothe autoclave. This value spreads from 70 to 90 vol% as listed in Table 3. A high fillinggenerates a higher internal pressure but provides a greater amount of solvent to dissolvemore mineralizer and consequently more ZnO species. We worked with 65e80 vol% in caseof our small autoclaves (16 mm inner diameter and 200 mm inner length) to achieve apprecia-ble growth rates.


52 (2006) 280e335

A closed Pt inner container [41,52] was found essential to prevent corrosion of the autoclaveinner walls due to the effect of the relatively higher basicity of the growth solution than thatused for artificial quartz [40e42,46,48,52,53]. In many former works Ag was used[44,45,47,61], which turned out to be less resistive against corrosion than Pt.

In our work we used a Pt inner crucible throughout [52]. The volume between autoclave andPt inner container was filled with a suitable amount of distilled water for pressure balancing toprevent the Pt inner container from serious deformation. The Pt inner container is filled up to80 vol% filling with the already prepared homogeneous solution of water, mineralizer and ZnOfeedstock. Seed holder and baffle are inside the Pt inner container and are both made of Pt.Fig. 3 shows the photograph of the Pt seed holder used for growth processes in autoclavesof 50 mm inner diameter and about 1 m inner length. However, this type of seed holder hasthe disadvantage of a relatively large surface area in relation to the volume of the inner Pt cru-cible. Often, parasitic nucleation can be found after the experiment. A general rule is thereforeto decrease the size of the seed holder as much as possible to diminish the surface available forparasitic nucleation while the mechanical stability is still guaranteed. What we typically ob-serve for small size autoclaves (16 mm inner diameter and 200 mm inner length) is the negativeeffect of seed holder and crystal on the mass flow. Sometimes crystals are not well-shaped, i.e.some facets were not fully developed due to effective shading and localized turbulence. Thisproblem can be reduced by reducing the number of seed crystals; consequently, the seed holderwould be smaller as well.

Fig. 3. Platinum seed holder as used for the autoclave of 50 mm inner diameter. The largest ZnO crystal (center) is about

25 mm across c-plane.


The above-mentioned trend gets less severe with larger autoclave size as shown in Fig. 3.The larger crystal (center) is about 1 inch in diameter across (0001) plane and has developedexcellent facets like most of the other crystals from this experiment. The growth of 2 and 3inch size crystals required to increase the inner diameter to 200 mm (length 3 m) and300 mm, respectively. Fig. 4 illustrates the result from one growth run. Almost 100 specimens2 inch in size have been grown on mainly (0001) and ð1010Þ oriented seeds. All crystals de-veloped excellent facets and the average crystal thickness was about 1 cm in the h0001idirection.

A ZnO crystal 3 inch in diameter and a wafer processed from it are shown in Fig. 5a and b,respectively. The crystal is about 1 cm in thickness. The light yellow coloration is due to im-purities from the ð0001Þ face, see Section 3.2.3 on impurity analysis. To date, these are the larg-est ZnO single crystals grown by the hydrothermal method.

The growth mechanism of ZnO under hydrothermal conditions has already been explainedby Laudise et al. [44], Khodakovsky and Elkin [67] and Demianets and Kostomarov [51]. Themain zinc species in the solution are ZnOOH�, Zn(OH)4

2�, ZnO22� and their concentrations de-

pend on the OH� concentration (i.e. pH value) and temperature. By far the major quantity ofthe OH� concentration derives from the basic mineralizers, but also the autoprotolysis of watercontributes according to:

2H2O4H3Oþ þOH� ð2Þ

The driving force to foster the growth process comes from the high concentration ofZnOOH�, Zn(OH)4

2�, ZnO22�in the solution, which builds-up the supersaturation necessary

to trigger nucleation on the ZnO seed crystal. Possible reactions leading to the growth ofZnO comprise:

ZnðOHÞ2�4 /ZnðOHÞ2þ 2OH� ð3Þ

and

ZnðOHÞ2/2Hþ þZnO2�2 : ð4Þ

Fig. 4. Two-inch size ZnO crystals produced during one growth run.


52 (2006) 280e335

From a comparison of the constant of reaction K:

K ¼ Kp=Kr ð5Þ

with Kp and Kr the constant of reaction of products and reactants, respectively, we get:

K ¼ ½Hþ�2½ZnO2�2 �=½ZnðOHÞ2� ð6Þ

and

½ZnO2�2 �=½ZnðOHÞ2� ¼ Kr½OH��2=K2

w; ð7Þ

with Kw the ionic product of water, it was shown that an increase of OH� in the solution yieldsan increase of ZnO2

2� [51]. The growth on the (0001) and ð0001Þ, i.e. Zn- and O-terminatedfaces or Znþ (surface) and O2þ (surface), respectively, likely involve:

ZnþðsurfaceÞ þZnO2�2 /2ZnOðcrystalÞ; ð0001Þ ð8Þ

O2�ðsurfaceÞ þZnO2�2 þ 2H2O/4OH� þZnOðcrystalÞ; ð0001Þ: ð9Þ

Znþ (surface) and O2þ (surface) describe those species that are provided by the crystal surface.The source for ZnO2

2� and H2O is the hydrothermal solution itself.In summary of reactions (2)e(9), during a ZnO growth experiment one might have the fol-

lowing flow of the Zn-containing species from the precursor to the crystal:

ZnOðprecursorÞ þ 4OH�/ZnðOHÞ2�4 þO2�/ZnðOHÞ2þ 2OH�

/ZnO�22 þ 2Hþ/ZnOðcrystalÞ þ 2e� ð10Þ

The growth speed is two times faster on the (0001) face than on the ð0001Þ face, and hasbeen explained by the ratio of formed ZnO units, ZnO (crystal), as obtained from formulae(8) and (9) [51]. This is close to the reported values of 2e3 [41,43,46,48,52].

Similarly, the growth speed is about two times faster on the (0001) face than on the ð1011Þface [43].

The presence of Liþ slows down the growth speed in the h0001i direction and slightly in-creases it for the h1100i direction [44,51]. This is considered to be related to a decreased

Fig. 5. View down c-axis (a) of a 3-inch ZnO crystal, (b) 3 inch (0001) ZnO wafer with CMP finish. Scale bar is 80 mm.


positive surface charge that lowers the probability of incorporating Zn-containing negatively-charged species. A high growth rate of 2.053 and 1.097 mm/day for the (0001) face andð0001Þ face, respectively, was reported by Demianets and Kostomarov [51], see Table 3.KOH was solely used as mineralizer but is known to reduce the crystal quality [38]. However,already from the old work of Laudise et al. [44] it is known that Liþ is needed to reduce thenumber of crystal defects (see also Section 3.2.1). This was confirmed by our work. Wefound that a molar ratio of 3:1 of KOH:LiOH delivers the best quality, i.e. highly transparentcrystals with a very small X-ray rocking curve full-width half-maximum (FWHM, seeSection 3.2.3).

The above discussion shows the general dilemma with mineralizers: they are indispensablein establishing an effective solubility of the precursor. On the other hand, they are unwelcomefor the process yield in crystal growth. This is intrinsic to the hydrothermal growth of ZnO andcan only be compromised as shown above.

Typical growth rates of high-quality ZnO crystals occur in the range up to a maximum of0.3 mm/day (Table 3). This is comparatively slow compared with the pressurized melt growthof ZnO, where up to 10 mm h�1 can be achieved [7]. Two-inch size (0001) ZnO produced byseeded chemical vapor transport (SCVT) grew at growth rates of <70 mm h�1 [67]. Neverthe-less, the high crystallinity and throughput in large autoclaves clearly indicates the favourablecommercial potential of hydrothermal growth technology. Thus, the use of large-size autoclavesfor ZnO like those used for the growth of quartz (Fig. 2, Table 2) would produce sufficient ZnOwafers for a variety of applications.

3.2.3. Characteristics of the material

3.2.3.1. Crystallinity, defects. The quality of (0001) substrates was investigated by rockingcurve measurements using the (0002) reflection. An RINT-2000 (Rigaku) diffractometer withCuKa radiation was employed in combination with a four crystal Ge (220) channel monochro-mator, beam divergence 12 arcsec, scan speed 0.01� min�1, step width 10�4�. The FWHMranged between 19 and 30 arcsec after chemicalemechanical polishing (CMP). This is betterthan values reported for ZnO wafer machined from pressure-melt grown ZnO, 49 arcsec[7,69], and comparable to CVT ZnO, FWHM around 30 arcsec [70]. Other groups reportedFWHMs of 43 and 37 arcsec for polished (0001) and ð0001Þ surfaces of hydrothermal ZnO,respectively [48], and 57 arcsec [49] for slightly N-doped hydrothermal ZnO. Contact-modeAFM measurements taken under air conditions reveal a root mean square (RMS) roughnessof 0.285 nm (Fig. 6a) and 0.155 nm (Fig. 6b). The X-ray reciprocal space map using the(0002) symmetric reflection (Fig. 7) shows a highly-symmetric single peak with the FWHMfrom the u scan of 15 arcsec. This value was confirmed by Wenisch et al. [71] who reportedabout 12e15 arcsec by triple-axis ue2q scans.

X-ray topography (CuKa radiation, 40 kV, 10 mA, detected by a film IX80; BergeBarrettgeometry) was measured on a (0001) ZnO wafer 2 inch in size and 500 mm in thickness.The XRC FWHM from the (0002) reflection of the sample was 19 arcsec. Two thousand fivehundred scans have been assembled to yield the contrast-enhanced image of Fig. 8. The(114) reflection at 2q¼ 98.6� and u¼ 49.3� was employed. The wafer appears very homoge-neous over the entire area. Slight contrast effects are seen, which is presumed to be due toslightly different lattice parameters caused by fluctuations in the impurity concentration or stoi-chiometry [46]. The lower part in Fig. 8 shows a few lines with a surface of apparently irregularcrystallographic orientation. Identification of the defect is thus impossible and more work is


still required. Croxall et al. [46] reported on dislocations as the principal imperfections in theirhydrothermal ZnO crystals. They lie in the basal plane and run parallel to or 20� inclined to thegrowth direction of h1010i. The point defects at the interface of the seed-grown crystal initiateddislocations.

The etch pit density (EPD) was determined for (0001) and ð0001Þ wafers. Fig. 9 shows theresult from etching with an aqueous solution of concentrated H3PO4 for 5 min at 25 �C. Theetching behavior is strikingly different as displayed by the shape of the etch figures. Whilethose on the (0001) face clearly exhibit a 6-fold axis with pyramidal facets, those on theð0001Þ face are less facetted. The EPD was about 300 cm�2 after chemicalemechanical polish-ing and was further lowered to less than 80 cm�2 by annealing [53]. The etch rate in the an-nealed wafers was low, which was related to the improved crystallinity. An aqueous solutionof 0.7% HCl was applied for 5 min at 60 �C.

Sakagami et al. [40] report an EPD of 100 and 103 cm�2 on (0001) and ð0001Þ faces, respec-tively, using an aqueous solution of H3PO4. Dislocations were responsible for the higher EPDon the ð0001Þ face.

Fig. 6. AFM images of ZnO wafers with different surface finish: (a) (0001) face after CMP, (b) ð0001Þ face after CMP,

(c) (0001) after annealing at 1050 �C and (d) (0001) after annealing at 1100 �C. Annealing time was 120 min; oxygen

flow rate was 200 ml min�1.

52 (2006) 280e335


Etching with 1% HCl yields an etch rate of 10 mm min�1 according to another report [72]and a mixture of 1 ml:1 ml:10 ml of conc. H3PO4:conc. acetic acid:H2O etched at 1.5 mmmin�1. The use of 30% HNO3 aqueous solution produces hexagonal pyramids on the (0001)face of mechanically polished ZnO [73].

3.2.3.2. Impurities. The incorporation of impurities and non-radiative recombination centersstrongly depends on growth sectors [41,42,44,46,48], which are characteristic of the

Fig. 7. The (002) reflection from the X-ray reciprocal space mapping of a (0001) ZnO wafer.

Fig. 8. Transmission X-ray topography of a polished high-quality ZnO wafer of 2 inch in size. The image is contrast

enhanced.


hydrothermal growth of ZnO. Growth sectors are formed due to differences in growth rate andmechanism for faces with different crystallographic orientations (see also Section 3.2.2).Fig. 10 is the result from the observation of a (0001) wafer by fluorescence microscopy(lexc¼ 365 nm) showing different intensities from the pyramidal, prismatic and basal sector.The intensities decrease from the prismatic to the pyramidal to the basal sector, which alsowas observed in the broad emission band at <2.8 eV in spectra derived by cathodolumines-cence (CL) by Mass et al. [42]. It was pointed out that samples cut from the crystal volumedirectly above the (0001) seed did not show other than that found in the low-impurity basalsector. Results from XRC measurements on 2 inch wafers using the (0002) reflection supportthese findings. The FWHM is at 18e22 and �50 arcsec for that part from only the basal sectorand for both prismatic and pyramidal sectors, respectively [63].

Consequently, this is used for the production of large-size ZnO wafers by Tokyo Denpa(TEW, Fig. 11). Little random strain was observed under crossed Nicols, which also supportsthe result from X-ray topography that the homogeneity of the wafer is very high.

Fig. 9. Etch pits on the (0001) and ð0001Þ surface of a polished ZnO wafer.

Fig. 10. Different growth sectors as revealed by fluorescence microscopy using excitation with lexc¼ 365 nm. The im-

age shows the view through a (0001) wafer displaying pyramidal, prismatic and basal sector.

52 (2006) 280e335


Secondary ion mass spectrometry (SIMS) was measured to study the impurity distribution inthe depth of a wafer (Fig. 12). The primary beam species was Csþ (5 kV, 350 nA); sputter speed120e150 nm min�1. As can be seen from Fig. 12 the impurity levels remain constant with in-creasing scan depth. The strong increase in intensity for all impurities measured over the last<100 nm is knowingly an artifact of SIMS.

However, a more sensitive investigation on impurity concentrations in 2 inch ZnO fromTEW has been made using inductively coupled plasma mass spectrometry (ICP-MS) [52,53].In Fig. 13 is shown the concentration of Fe, Al, Li, K and its position dependence. A negative

Fig. 11. Little random strain is revealed by optical observation under crossed Nicols of a 2-inch ZnO wafer after CMP.

Fig. 12. SIMS depth scans of impurities in a ZnO wafer with CMP surface finish.


52 (2006) 280e335

wafer number (�3 to �1) indicates those specimens grown and cut from the ð0001Þ face andpositive ones (1e4) refer to specimens from the (0001) face of the seed crystal. Increasingdistance from the ð0001Þ face of the seed generates less impurity in the grown crystal. Thisis particularly obvious for Li with 12 and <1 ppm for wafer numbers �1 and �3, respectively.The concentration of K remains unchanged for both faces at <0.3 ppm. Both Fe and Al showhigher concentrations in wafers grown on the ð0001Þ face of the seed crystal [52,53]. Theconcentration numbers are <11 and <1 ppm for Fe and <8 and <0.5 ppm for Al for theð0001Þ face and (0001) face, respectively. Mass et al. [42] report about 1e10 ppm of Li, Kand Na not above 1e2 ppm and Al, Fe, Si, C from 1 to 10 ppm.

The result from the examination of the ZnO precursor chunks and 2 inch ZnO crystals byglow discharge mass spectrometry (GDMS) is summarized in Table 4. Argon was used asa discharge gas. The result for the precursor was measured at the end of the year 2001,that for the crystals in early 2005. The purity of the precursor has been improved duringthis period. The Pt concentration in the crystals, however, is clearly assigned to the inner linermade of Pt.

A ZnO crystal grown by CVT using H2 and N2 as carrier gas [6] incorporates0.03e0.1 ppm K and 0.2e1 ppm Na as recorded by AAS. Fe and Al were found at concentra-tions below the detection limit of 1 ppm and 7 ppm, respectively. The group of Triboulet et al.[70] uses Cl2 and C as transport agents in CVT and have detected about 0.053% of Cl and0.05% of C when using graphitized ampoules. For comparison, pressure-melt grown ZnO[69] contains 4 ppm Fe, 8 ppm Pb and 2.5 ppm Cd.

Hydrogen-related defects in hydrothermal, as-grown ZnO (1017 cm�3 carrier concentrationat room temperature, see also Fig. 19a) have been studied by infrared absorption spectroscopy[74]. A number of lines are located in the vicinity of the characteristic OeH stretch local vi-brational modes and the line at 3577.3 cm�1 was related to a defect containing one OeHbond primarily aligned with the c-axis of the crystal. The presence of a substitutional Niatom in the defect was tentatively proposed.

The same absorption line of 3577.3 cm�1 at 12 K, and shifting to 3547 cm�1 at 300 K, wascharacterized by electron paramagnetic resonance (EPR) in a hydrothermal ZnO sample [75].This line, however, was assigned to an OH� ion located on an oxygen site adjacent to an Liþ

Fig. 13. Concentration analysis of Fe, Al, Li, K as detected by ICP-MS in seven wafers cut from one crystal.


ion on the zinc site. The concentration of the neutral complex Li�OH� was estimated to be7.6� 1017 cm�3. Charge compensation by singly ionized acceptors was concluded to play a ma-jor role for hydrogen in ZnO. A band at 3546 cm�1 at 300 K [64] has already been suspected tobe caused by hydrogen located in a bond-centered position between oxygen and zinc. This bandmight contain some unresolved components from OH� involving an acceptor on the zinc site.Furthermore, detection of Co2þ at 6005 cm�1 and Ni2þ at 4240 cm�1 was reported. Evidenceof Fe3þ and Mn2þ was also found in the same hydrothermal ZnO sample [75]. Photo-inducedEPR measurements revealed a signal due to neutral Li acceptors. The concentration of photo-induced neutral Li acceptors was approximately 1.3� 1015 cm�3.

Recently, Demianets et al. [38] reported on the modification of the crystal habit upon incor-poration of some bi- and tri-valent metal ions like Ni2þ, Cd2þ, Mn2þ, Co2þ, Fe2þ, Fe3þ andIn3þ. The ions Ni2þ, Cd2þ, Mn2þ, Co2þ and In3þ substitute for the Zn site in the crystal latticeand Ni2þ, Mn2þ, Fe2þ and Fe3þ were said to cause crystal coloration. In the case of In3þ theformation of a point defect of the type InZn

þ was found to be a shallow donor. This effect wasrecently used in combination with Li doping to achieve super-fast luminescence decay bydonoreacceptor pair recombination [97]. Due to the relatively low growth temperature underhydrothermal conditions, formation of InZn

þ LiZn0 type associates does not happen.

3.2.3.3. Thermal properties. The linear thermal expansion coefficient DlL/L, specific heat cp

and thermal diffusivity a of hydrothermal grown ZnO have been evaluated for samples

Table 4

Impurity analysis of 2-inch size ZnO crystals and ZnO precursor chunks

Element Concentration (ppm wt) Remarks

Crystal Chunks

Li 1.3 19 From mineralizer

Na 0.11 0.13

K <0.1 6.5 From mineralizer

Rb <0.01 e

Cs <0.05 eMg 0.04 0.25

Ca <0.1 e

Ti <0.005 e

V <0.005 <0.005

Cr <0.05 0.15

Mn <0.05 e

Fe 0.06 3.7

Co <0.01 eNi <0.01 <0.01

Cu <0.01 <0.01

Pb 0.07 eCd <0.1 <0.03

Al 0.02 3.4

Ga <0.05 0.09

Si 0.18 1.9

Ge <0.05 <0.05

As <0.05 <0.05

Sb <0.5 e

Pt 7.1 <0.02 From Pt liner

Lanthanides <0.05 e Pm was not measured


52 (2006) 280e335

prepared from hydrothermal ZnO grown by TEW and for ZnO powder in case of cp; the thermalconductivity k was calculated from the measurement of a.

The thermal expansion DL/L was measured with a dilatometer (RIGAKU Thermo plus TMA8310) in comparison to the standard sample a-Al2O3 (sapphire). The sample size was4� 4� 19 mm3 for both measurements in the h0001i and the h1120i directions. Fig. 14 showsDL/L versus temperature for the temperature interval 20 �C� T� 1100 �C. The following lin-ear thermal expansion coefficients have been calculated: 6.06� 10�6 K�1 and 4.16� 10�6 K�1

at 25 �C and 10.4� 10�6 K�1 and 4.87� 10�6 K�1 at 1000 �C for the h1120i and h0001i di-rections, respectively.

Cermet Inc. [76] gives the value of DL/L¼ 2.9� 10�6 K�1 (no temperature range specified,but obviously for h0001i direction at RT) for their wafers were prepared from melt-grown ZnOingots.

Based on first principle theory, Reeber and Wang [77] have calculated DL/L for the h1120iand h0001i direction, DL/L¼ 4.867� 10�6 K�1 and 2.911� 10�6 K�1 (298 K) and9.403� 10�6 K�1 and 5.439� 10�6 K�1 (1400 K), respectively. These results are quite consis-tent with our measured values.

For comparison with hexagonal GaN, the linear thermal expansion coefficient at RT alongthe a-axis and the c-axis is 5� 10�6 K�1 and 4.5� 10�6 K�1, respectively [78]. Particularlythe difference in h0001i direction between ZnO and GaN is quite similar.

The specific heat of ZnO powder (99.999% purity) has been measured using a PerkinElmer DSC-7. Fig. 15 shows the temperature dependence of cp for the range 20e300 �C. Thespecific heat of ZnO is 0.492 and 0.603 J g�1 K�1 at 20 and 300 �C, respectively. Melt-grownZnO from Cermet [76] is quoted as 0.523 J g�1 K�1 (apparently at a temperature related to theRT region e no temperature was specified in the reference).

The thermal diffusivity was measured using the laser flash method as described by Wasedaet al. [79]. A (0001) and a ð1120Þ oriented, polished and Au-coated ZnO disks of 1 mm inthickness and 10 mm in diameter were exposed to a laser pulse (l¼ 694 nm) and the emittedradiation was measured by an IR sensor. The following formula was used to calculate a [80]:

Fig. 14. Thermal expansion along a- and c-axis of hydrothermal ZnO from TEW for the temperature range of

20e1000 �C.


a¼ 1:38L2=P2t1=2; ð11Þ

where L is the thickness of the sample and t1/2 the time required for the temperature response toreach 50% of its maximum value. At 25 �C we obtain 1.76� 10�5 and 2.05� 10�5 m2 s�1 forthe a-axis and the c-axis, respectively.

The thermal conductivity k can be calculated accordingly to:

k¼ arcp; ð12Þ

where r is the density of ZnO. The values for k are 49.1 and 57.2 W m�1 K�1 for the a- andc-axis, respectively.

Melt-grown ZnO from Cermet [76] is referenced to as k¼ 130 W m�1 K�1 (no temperaturespecified), respectively.

For comparison k of sapphire, which is widely used as a substrate material in GaN epitaxy, is46.06 W K�1 m�1 at 273 K [82] which is a little lower and will be even lower at 298 K.

3.2.3.4. Optical properties. The optical transmittance of (0001) ZnO wafers from TEW (pol-ished sample of 0.5 mm thickness, no visible scratches on surface) [63], hydrothermal ZnOfrom SPC Goodwill (GW; polished sample of 0.5 mm thickness, no visible scratches on sur-face; purchased in the year of 2004) and a thin, 1% In-doped hydrothermal ZnO crystal (plateletof 3 mm in diameter and 0.2 in thickness with surfaces as grown) grown by TEW is shown inFig. 16. The sample from TEW shows a transmittance of 80% at l¼ 410 nm and 87% atl¼ 700 nm. The sample from GW, which appeared a touch more yellowish than the TEW spec-imen, shows a slightly higher absorption in the range 390 nm� l� 500 nm. This could be as-signed to a slightly higher impurity level in this GW sample. In-doping lowers the transmittanceto about 60%. The measurement has been performed on about 200 mm thin crystals. The asgrown (0001) surfaces were used and loss due to surface scattering might marginally distortthe result.

The refractive index for the ordinary (no) and extraordinary (ne) beam of hydrothermal ZnOfrom TEW and GW is shown in Fig. 17. The samples are the same as for above measurement of

Fig. 15. Heat capacity of ZnO powder for the temperature range of 20e300 �C.


52 (2006) 280e335

optical transmittance. The laser beam was coupled into the sample by a rutile prism. In contrastto the optical transmittance identical values were obtained for each sample, which can be fittedby a second-order exponential decay.

The photoluminescence (PL) signal was obtained by excitation using a continuous-waveHeeCd laser (lexc¼ 325 nm, Pout¼ 1.6 mW) at 4 K. The signal was detected by a CCD camera(Princeton Instruments Inc.) after dispersion with a 30 cm triple grating monochromator. ThePL spectrum of a TEW ZnO sample is shown in Fig. 18. The broad emission band from 1.7to 2.8 eV peaks around 2.3 eV. The nature of this broad emission involves donoreacceptorpair recombination due to Li [83,84]. Copper impurities at levels lower 1 ppm have been as-signed to emissions around 2.4 eV [85e87] and Fe3þ shows well-defined lines at 1.7e1.8 eV [87]. The latter has not been observed in our samples.

Fig. 16. Optical transmittance of 0.5 mm thick hydrothermal grown (0001) ZnO wafers from TEW and GW. The 1%

In-doped ZnO sample was a 0.1 mm thin crystal with natural {0001} surfaces.

Fig. 17. Refractive indices for the ordinary (no) and extraordinary (ne) beam of light of hydrothermal grown ZnO from

TEW and GW.


The luminescence in the region near the band edge (see inset of Fig. 18) is clearly pro-nounced with peaks at 3.233, 3.307 (with a shoulder at 3.33 eV), 3.356, and around 3.37 eV,which can be assigned due to a donoreacceptor pair (DAP) transition, two-electron satellitetransitions (TES), neutral acceptor bound exciton (I9) recombination, and a free A-exciton re-combination (FEA), respectively [5,88,89]. At lower energy the phonon replicas (LO phononenergy¼ 72 meV) appear: �1LO at 3.161 eV and �2LO at 3.09 eV. Hydrothermal ZnO re-ported in other literature reveal similar PL and CL spectra [42,49,53]. This hints to comparableimpurity and defect concentrations in hydrothermal ZnO grown by the various groups.

Our picture of PL from hydrothermal ZnO over the entire energy range of about 2e3.5 eV isquite consistent with samples grown by SCVT by Eagle Picher [5,87]. Interestingly, the broadyellow band peaking around 2.3 eV also appears in the SCVT sample although the growth con-ditions are very different. This might lead to the assumption that the levels of active impuritiesin our hydrothermal ZnO are comparable to those of SCVT ZnO. Hence, this suggests that thehydrothermal growth technique is an economic way to yield a large quantity of high-quality,large-sized ZnO crystals (Table 3).

3.2.3.5. Electrical properties. There is some concern of the impact of surface properties on thereliability of results from measurements of electrical properties using surface contacts. It isknown that the ZnO surface does adsorb molecules, such as CO2, CO, O2, H2, etc., as revealedin the work by Gopel and Esser [90,91]. The formation of a highly conductive layer on the sur-face of highly resistive ZnO crystals was reported by Markevich et al. [92] and Schmidt et al.[93]. This was explained by adsorbed oxygen atoms on the surface of a ZnO crystal, which ledto the capture of electrons. This results in a negative surface charge and a depletion layer withreduced conductivity. The thickness of this layer on undoped ZnO was given as <1 mm [92].

Temperature-dependent Hall-effect technique with Van der Pauw geometry was used to ex-amine the carrier concentration (N, Fig. 19a), carrier mobility (mH, Fig. 19b), and electrical re-sistivity (R, Fig. 19c) from a 10� 10 mm2 specimen cut from a high-quality TEW ZnO crystal.The surface was polished to an RMS roughness of 0.2 nm. The Ti/Au contacts were producedby thermal evaporation. A melt-grown sample was measured for comparison.

Fig. 18. PL spectrums from a TEW wafer for the range 1.7e3.6 eV and 3.18e3.43 eV (inset) at a temperature of 4 K.


52 (2006) 280e335

Fig. 19. Comparison of the temperature dependence of (a) carrier concentration, (b) carrier mobility and (c) resistivity of

hydrothermal TEW ZnO and pressure-melt grown ZnO (all measured at p z 0.1 Pa).


The carrier concentration is very much lower in the hydrothermal sample than in the meltgrown one and decreased with increasing 1/T from N¼ 2� 1016 cm�3 at 500 K (103/T¼ 20)to N¼ 4� 1013 at 100 K. Polyakov and co-worker [94] obtained very similar results ofN¼ 1.3e4.6� 1013 and N¼ 6.4� 1011 at 300 K and 77 K, respectively, for measurementson four samples of TD ZnO as purchased. In both cases, the N value is clearly the effect oflow-impurity, point defect, and dislocation concentrations, quite in good agreement with ourresults from XRD and impurity analysis. The above species are known to produce electri-cally-active centers [95]. The slight hysteresis slope at 103/T< 4 was obtained from the mea-surement during heating up and cooling down and is likely the effect related to surfaceconductivity [92,93]. For surface conductivity see also the discussion in Section 3.2.4.2.SCVT grown ZnO showed a slightly lower carrier concentration at 103/T¼ 20, N¼ 3� 1014

[95]. In the same paper, annealing the SCVT sample at 950 �C in an He atmosphere lowersN to about 1014, which was ascribed to a reduction of most of the shallowest center, or at leastthe result of the Fermi level dropping below the energy of this center.

The Hall mobility peaks at 100 K, mH¼ 530 cm2 V�1 s�1 and drops down to about40 cm2 V�1 s�1 at 580 K. The higher mobility in comparison to the melt-grown sample, thatpeaks at 480 cm2 V�1 s�1 at 80 K and 430 cm2 V�1 s�1 at 100 K, is due to the lower impurityconcentration as indicated by the lower carrier concentration (Fig. 19a). A formerly hydrother-mal grown sample [95] had already showed very similar results of mH for the measured temper-ature range of 200e400 K. The mH lowered from about 300 to 100 cm2 V�1 s�1. SCVT grownZnO shows higher mH up to almost 2000 cm2 V�1 s�1 at 40 K. Compared to GaN, one wouldfind a lower mobility because of a higher effective mass and larger optical phonon scatteringparameter [95].

The electrical resistivity of the hydrothermal ZnO sample (Fig. 19c) is about two orders ofmagnitude higher than the sample grown from the melt [76], with the minimum of 20 U cm at60 K and 0.1 U cm at 200 K, respectively. Hydrothermal ZnO from GW shows a higher elec-trical resistivity of 500e1000 U cm [81] than the TEW material. One could speculate that theLi concentration must be higher in the GW crystal. However, no impurity data were available.Other results on TEW ZnO [94] reported on a large variation of R between 96 and5� 105 U cm, which was speculated to come from the Li concentration in the samples. It ispossible that different growth sectors were present in the specimens and therefore Li was incor-porated in quite different concentrations there.

The uniformity of R, N and mH over a 2 inch wafer from TEW was measured (Fig. 20) andgives following values for R¼ 380 U cm� 15%, N¼ 8� 1013 cm�3� 20%, mH¼ 200cm2 V�1 s�1� 10%, respectively [53]. These uniformities are quite satisfactory.

3.2.4. Machining of ZnO

3.2.4.1. Cutting and polishing. We used a diamond blade saw for cutting ZnO. The damageinduced by cutting is removed by a lapping process. A subsequent mechanical polishing is fol-lowed by the chemicalemechanical polishing (CMP), both serving for removal of the damagecaused by lapping. The polishing was done with a silica-based alkaline (KOH) slurry ofpH¼ 10.5 firstly, followed by a second step using a pH of 8e9. The grain size of the polishingpowder has been reduced successively from 10 to <2 mm. The typical FWHM of the XRCusing (0002) reflection is between 25 and 40 arcsec for both (0001) and ð0001Þ face ofa scratch-free ZnO wafer.


52 (2006) 280e335

Generally, mechanical machining causes surface damage in the ZnO wafer. The surfacedamage in hydrothermal ZnO substrates as supplied has already been reported by Wenischet al. [71] using double-axis (0002) HRXRD ue2q rocking curve measurements. A shoulderon the low angle side of the main peak was related to a larger lattice constant. The same groupof Yao et al. [96] later confirmed by (0002) HRXRD ue2q scans that a damaged surface layerunder tensile strain exists up to 15 mm deep into the volume of a CMP ZnO wafer. In the lowtemperature (10 K) PL spectrum relative intense first order phonon-assisted luminescence wasclearly found at 3.294 eV and a highly intense emission from the ionized donor bound excitonat 3.370 eV was revealed, both of which were assigned to crystal defects.

We found evidence for damage in the surface layer by comparing the results obtained byradioluminescence (RL) measurements of a ZnO hydrothermal substrate and an LPE-grownZnO film processed on this substrate (Fig. 38a) [18,97]. The LPE film is considered asa highly-crystalline, naturally grown ZnO face. The low temperature (80 K) RL measurementshows a five times higher intensity of the excitonic emission from the LPE film. The same emis-sion from the machined wafer, by contrast, almost disappears at room temperature but is clearlyvisible for the film. Moreover, the ratio of the intensity from the excitonic and green-yellowemission was about 0.1 and 2 for LPE film and the substrate, respectively.

The damaged layer in ð0001Þ bulk crystals may be removed by etching with CF3COOH [96],however, surface roughening is then observed. This caused the milky appearance of the treatedsurface.

3.2.4.2. Effects from annealing. Generally, the annealing aims to improve the crystallinity ofthe machined ZnO wafers. Particularly important is the control of the surface morphology ofwafers to be employed in subsequent epitaxial growth. Here, monatomic steps are desirableto control the epitaxial growth process at an atomic level.

We used a fused silica tube to define atmosphere control. This was placed in a horizontalfurnace. An O2 flow of 200 ml min�1, at temperatures from 600 to 1100 �C and annealing timesbetween 10 and 900 min were applied. Neither temperature nor process duration affected the

Fig. 20. Uniformity of resistivity, carrier concentration and mobility across a 2 inch epi-ready ZnO wafer provided by

TEW.


surface morphology. We found that the annealing time needed to be decreased with increasingtemperatures in order to achieve comparable results.

Fig. 6c and d shows the results of annealing for 120 min at 1050 and 1100 �C. The1� 1 mm2 AFM scan on the (0001) face of a ZnO wafer cut with 0.5� miscut angle towardh1010i displays the beginning of the formation of monatomic steps (Fig. 6c). The RMS rough-ness decreased to 0.096 nm. Increasing the annealing temperature to 1100 �C for 120 minyields steps of 1 nm in height equivalent to 2 unit cells (lattice parameter c¼ 0.52066 nm)with the interstep distance varying from 120 to 150 nm (Fig. 6d). The RMS roughness was0.1713 nm as a result of the formed steps.

Fig. 21 is an SEM image taken from the ð1010Þ face. The steps of 50 nm in width appearedupon annealing at 1050 �C for 120 min and O2 flow of 200 ml min�1. Apparently, the formationof steps on the m-plane requires a little less thermal activation than on the c-plane. See also theAFM image of a (0001) face after annealing under similar conditions (Fig. 6c).

The evolution of the morphology of the ð0001Þ face with the annealing parameter temper-ature was reported by Ko et al. [98]. The samples were directly exposed to a high purity(99.999%) oxygen stream for 60 min. It was found that a 1000 �C annealing temperature is suf-ficient to improve the crystallinity as proven by the low (0002) rocking curve FWHM of17.2 arcsec. The RMS roughness was reduced from 3.1 nm (surface quality as supplied) to0.3 nm for the sample annealed at 1000 �C. Atomically flat surfaces with a step height andwidth of 0.5 and 280 nm, respectively, were obtained for the annealing temperature.

The annealing of the (0001) face at 1000 �C for 1e5 h in an oxygen atmosphere was re-ported by Cho et al. [96]. Steps were obtained by annealing for �1 h. They are aligned inthe h1010i direction. It was shown that the step height is a function of the annealing timeand rose linearly, i.e. 2 nm at 1 h to 10 nm for the 5 h anneal. Annealing for 3 h yielded thebest crystallinity as indicated by the FWHM from the rocking curves of (0002), (0004),(0006), and ð1011Þ reflections of 12.6, 12.24, 12.96, and 16.56 arcsec, respectively. Annealingfor longer than 3 h deteriorated the crystalline quality.

Fig. 21. SEM image of the m-plane showing steps of about 50 nm in width obtained by annealing at 1050 �C for

120 min; oxygen flow rate was 200 ml min�1.


52 (2006) 280e335

The above paper [96] also reports on annealing under N2, O2, and high vacuum conditions(�5� 10�9 Torr) at 950 �C for 4 h. Desorption or formation of pits at the surface occurs foreither an N2 or an O2 atmosphere, respectively. Clear steps were obtained by annealing underhigh vacuum giving an RMS roughness of 0.25 nm. Surface reconstruction of a 3� 3 patternwas observed by RHEED by annealing over 30 min. Annealing over 3 h slowly decreasedthe RHEED intensity. Using an oxygen RF plasma and applying a chamber pressure of5� 10�5 Torr did not lead to an observation of a reconstruction pattern.

It has to be noted here that results from annealing ZnO strongly depends on the experimentalconditions governed by the setup of the equipment. This is confirmed by the slightly differentobservations as shown above.

Fig. 22 shows that annealing may cause the formation of voids as observed on the surface ofa polished and annealed (0001) sample. The size of the void typically is less than 1 mm butsometimes extends to more than 10 mm and exhibits a layer consisting of sub-micrometersize needle crystals (region B in Fig. 22). It could be speculated that crystal defects likea high local impurity concentration might be the cause that initiates the formation of voids.However, no further experiments have been conducted so far. The surface surrounding the voidsappears with the typical step formation (region A of Fig. 22).

Scratches caused by improperly polishing do heal out in parts as demonstrated by Fig. 23.Hexagonal pits of some hundreds of nanometers and smaller irregular voids are visible in thetrace of the scratch (located by the darker bar in the center of Fig. 23). The voids in the upper

Fig. 22. The SEM image of the (0001) plane after annealing displays a large void. The bright part consists of recrystal-

lized ZnO (B). The surface around the void appears with typical step pattern (A).


part of the image do not relate to this polishing effect, but rather to the above-mentioned for-mation of voids.

The impact of annealing on electrical properties of the ZnO surface was recently reported bySchmidt et al. [93]. Upon annealing at �650 K in a vacuum of <0.1 Pa, a semi-insulating (i.e.highly resistive due to Li doping) specimen showed a high surface conductivity. This was ex-plained by the accumulation of electrons in the potential well near the surface. Thus, an elec-trically conducting layer was formed. Subsequent annealing in air recovered the originally highresistivity.

Fig. 24 shows the Al concentration profile of a specimen before and after annealing. TheSIMS data reveal a slight decrease in the Al concentration from 6� 1016 to 3� 1016 cm�3.

Fig. 23. SEM image of the (0001) plane of a ZnO wafer after annealing at 1100 �C shows the healing effect of anneal-

ing. Large hexagonal pits, smaller voids and steps were formed. The trace of the scratch induced by polishing has been

visualized by the darkened bar.

Fig. 24. SIMS depth scans demonstrating the effect of annealing on the Al concentration in a ZnO wafer.


52 (2006) 280e335

However, the Al concentration in the annealed sample slightly increases toward the surface andexceeds the value from the untreated sample in the volume up to 200 nm deep into the sample.Again, this might rather be related to the SIMS measurement than the true properties of thesample.

Polyakov et al. [94] have reported on the out-diffusion of Li at temperatures as low as500 �C. Mainly the surface near region was affected. This effect indeed might be useful totune the resistivity of the ZnO surface.

4. Solvothermal film growth of ZnO

4.1. Growth from aqueous solvents

The limited choice of solvents for ZnO has already been discussed in an earlier section ofthis paper and fully applies to the growth of ZnO films. In thin film technologies, however, weare faced with an additional challenge which is related to the substrate. The film growth processhas to be such as to avoid substrate corrosion, i.e. dissolution or any kind of ion exchange withthe solution should be minimized in order to obtain a defined interface between substrate andgrown film. This is often the limiting aspect in film growth technology. Typically, ZnO filmgrowth is based on chemical reactions to produce a ZnO solution which consists of a solvent,precursor and a complexing agent. The latter through pH (concentration of H3Oþ) control[39,89,100] is responsible for the crystal habit.

The majority of growth techniques are based on aqueous solvents including waterealcoholmixtures. Spin-coating is a non-aqueous technique using 2-methoxyethanol, which will beevaporated after the coating process. Similar to water, 2-methoxyethanol contains an OH�

group, which may be incorporated in the ZnO film. The potential of alkaline-metal chloridesas water-free solvents will be treated in detail in the following Section 4.2.

Common to all water-based growth techniques is the low process temperature, even<100 �C in air under atmospheric pressure. The removal of growth solution from the fabricatedfilm is trouble-free. The growth of ZnO films has been demonstrated by a variety of methods:chemical bath deposition (CBD), selective ionic-layer absorption and reaction (SILAR), elec-troless deposition (ED), liquid flow deposition (LFD). A comprehensive description is given byNiesen and De Guire [100]. CBD produces a solid film in a single immersion through control ofthe kinetics of the formation of ZnO. SILAR is characterized by the use of alternating aqueoussolutions containing a zinc salt and a hydrolyzing solution (water or water/ammonia). ED usesoxidizing or hydrolyzing agents to form ZnO and is similar to CBD. Films grown are mostlywithout crystallographic texture. LFD involves the flow of the solution past the substrate ata controlled rate and reactants are replenished continuously. This process can be maintainedover longer time span to deliver thick films.

The above methods are used in ceramics technology but might be adaptable to the growth ofsingle-crystalline film by choice of a proper single-crystalline substrate. Table 5 gives an over-view over achievements in the growth of poly- and single-crystalline ZnO film from aqueoussolvents including 2-methoxyethanol. The application of glass-like substrates, plastics or sili-con wafers causes the growth of islands and polycrystalline ZnO films composed of grainsor columns.

The best crystallinity was obtained by using either single-crystalline (0001) ScAlMgO4

(SCAM) or a (111) MgAl2O4 substrate. SCAM is isostructural with YbFe2O4, having a spacegroup R3m with lattice constants a¼ 3.246 A and c¼ 25.195 A. The lattice misfit in the a plane

Table 5

Result Refs.

Method; film thickness; growth rate;

film morphology

Growth of polycrystalline films with

hexagonal grains z 0.3 mm;

�0.6 mm h�1

[101]

Boron-containing films up to 1 mm;

0.42e0.58 mm h�1[102]

Nanocolumnar films up to 3 mm;

growth time 1 h

[103]

Post-growth annealing at above 500 �C;

100 mm thickness (several dippings

required)

[104,105]

Spin-coating at 5000 rpm for 30 s;

post-growth annealing at max. 850 �C;

13e25 nm thick films.

[106,107]

Spin-coating at 3000 rpm for 30 s;

post-growth annealing at max. 700 �C;

nano-crystalline films mainly oriented

with c-axis perpendicular to substrate

surface, grain size� 80 nm.

[108]

Film thickness z 15 mm; Island growth

and coalescence

[19,109]

Epitaxial lateral overgrowth; (1) thin

ZnO film and (2) growth from window

region of film (1).

[110]

Microwave activation. Films of

rods mainly oriented with c-axis

perpendicular to substrate surface,

few mm thick.

[111]

31

2D

.E

hrentraut

etal.

/P

rogressin

Crystal

Grow

thand

Characterization

ofM

aterials52

(2006)280e

335

Growth conditions of solvothermal ZnO films

Growth conditions Growth vessel

Solvent Precursor/complexant T (�C) pH Substrate

H2O 0.05 mol l�1 Zn(NO3)2þ� 0.15 mol l�1 DMABb

50 6.2 Corning 7059 glass;

surface catalyzed with

SnCl2 and PdCl2

n.ra

0.1 mol l�1 Zn(NO3)2þ� 0.1 mol l�1 DMABb

60 Corning 1737 glass;

surface catalyzed with

SnCl2 and PdCl2H2O 0.005 mol l�1 Zinc carboxylate

salts (acetate, formate)þ0.005 mol l�1 HMTd

90 5.0 TOFc glass; Au coated Open system, volume

100 cm3

2-Methoxyethanol 0.75 M Zn(OAc)2�2H2OeþMEAf

RT n.r. Silica glass Open system

2-Methoxyethanol 0.75 M solution of Zn(OAc)2�2H2OeþMEAf

80 / (0001) ScAlMgO4;

(0001) a-Al2O3

/

2-Methoxyethanol 0.35 M Zn(OAc)2�2H2Oe, MEAf

RT / Corning 1737 glass /

H2O 0.025e0.1 mol l�1 Zn(NO3)2�6H2OþNH4OH

150e200 7e11 (111) MgAl2O4 Teflon-lined stainless

steel autoclave,

vol. 45 ml

H2O (1) Zn(NO3)2� 6H2OþNH4OHþNH3;

80 7.5 (111) MgAl2O4

(2) Zn(NO3)2� 6H2OþNH3þNa(cit)� nH2Og

90 10.9

H2O Zn(NO3)2, Zn(OAc)2�2H2Oe, urea, HMTd

�100 n.r. Glass covered by

F:SnO2

Pyrex reactor, open

system, vol. 80 ml

H2O 91 mM Zn(OAc)2� 2H2Oe,

4.5 mM NaOH

RT (?) 6.7 (100) Si, amine-

derivatized

Open system Solegel synthesis; Island growth,

substrate coverage 70%, 3e5 nm

thickness

[112]

0 ml

ate

Chemical spray pyrolysis of In:ZnO,

nano-crystalline films, grain

size� 34.7 nm, film thickness of

0.54 mm. Preferential c-axis orientation

[113]

ith

,

Aligned nanowires, film thickness

3 mm, growth time 10e20 h.

[114]

g Aligned nanowires, film thickness

3 mm, growth time 0.5e6 h.

[115]

31

3D

.E

hren

trautet

al./

Progress

inC

rystalG

rowth

andC

haracterizationof

Materials

52(2006)

280e335

H2Oþ ethanol 0.6 M Zn(OAc)2�2H2Oeþ InCl3

400� 5 n.r. Soda lime glass Open system, 20

solution, spray r

of 10 ml min�1

H2O 0.1 M Zn(NO3)2�xH2OþHMTd

95 n.r. ZnO-coated Si, PETh,

a-Al2O3

Reaction flask w

reflux condenser

40 ml solution

H2O 0.025 M Zn(NO3)2�xH2Oþ 0.025 M HMTd or DTi

90 n.r. ZnO-coated (100) Si,

PDMSjOpen crystallizin

dish

a n.r. e Not reported.b DMAB e dimethylamineborane.c TOF e fluorine-doped tin oxide.d HMT e hexamethylenetetramine.e Zn(OAc)2� 2H2O e zinc acetate.f MEA e monoethanolamine.g Na(cit)� nH2O e sodium citrate.h PET e polyethyleneterephthalate.i DT e diethylenetriamine.j PDMS e polydimethylsiloxane.

314 D. Ehrentraut et al. / Progress in Crystal Growth and Characterization of Materials

of (0001) SCAM to (0001) ZnO is 0.09% [117]. SCAM is composed of alternating layers ofwurtzite-type (0001) (Mg, Al)Ox and rock-salt type (111) ScOy, which is the reason for theexcellent (0001) cleavage. Recently, the first vapor-grown blue LED has been demonstratedusing (0001) SCAM as substrate [3]. It has also been used as substrate in ZnO film growthby spin-coating CBD and successive thermal annealing at a maximum of 850 �C [106]. The(0001) ZnO films were epitaxially grown in-plane with the (0001) SCAM substrate as revealedby XRD and HRTEM. The crystallographic orientation relationship between ZnO and SCAM isð0001ÞZnO ð0001ÞSCAM

�� and ½2110�ZnOk½2110�SCAM.The following remarkable results have been achieved by the group of Lange from UCSB.

Thick epitaxial films of about 15 mm were grown on (111) MgAl2O4 (spinel, space groupFd3m) at �200 �C [19]. The growth solution was composed of different concentrations ofZn(NO3)2� 6H2O, NH4OH, NH3, and Na-citrate (Na3C6H5O7� 2H2O) at pH between 7 and11 [19,110,111]. Fig. 25 shows the SEM image of a sample grown by two-cycle lateral epitaxialovergrowth (LEO), with the second film not completely coalesced [110]. This second film dis-plays no etch pits on the planar surface, which is the result of growth from window regionslying above the wing regions of the first film. A photo resist (PR) channel stamping techniquehas been used to produce growth windows [118]. The crystallographic relationship betweenZnO and MgAl2O4 was determined by XRD. The out-of-plane orientation ish111ispinel h0001iZnO

�� , the in-plane orientation is indicated to ½1120�ZnOk½112�spinel and½0110�ZnOk½110�spinel, respectively. The reduction of the lattice mismatch of ZnO on (111)MgAl2O4 from 13.6% to �1.6% was explained by a 30� rotation of oxygen planes in thewurtzite structure, which converts the Mg and Al sites in the spinel structure to tetrahedrallycoordinated Zn sites in ZnO. Hþ ions are necessary for charge balance at the interface betweenthe ZnO and MgAl2O4. The same group [110] recently reported on lateral epitaxial overgrowthby applying a stamped mask to a previously grown ZnO film. The growth of coalesced ZnOfilms was achieved at temperatures� 90 �C and required the addition of sodium citrate tothe solution. The film growth started with island growth which eventually coalesced to a closedfilm [19].

Fig. 25. SEM image showing lateral epitaxial overgrowth of ZnO at 90 �C (courtesy of Jin Hyeok Kim, David Andeen,

and F.F. Lange).

52 (2006) 280e335


52 (2006) 280e335

It was shown by Chaparro [119] that Zn(OH)2 is the predominant zinc species in aqueoussolutions containing ammonia as a complexing agent at pH values between 7.6 and 11.4. Am-monia plays an outstanding role as a complexing agent, since it forms metastable complex ionsthat can result in the complete dissolution of zinc, establish an appropriate basic pH that decom-poses the oxygen delivering precursor, and also possesses a sufficiently high stability in thesolution.

4.2. Liquid phase epitaxy from chloride solution

Liquid phase epitaxy (LPE) from chloride solution has recently been proven successful forthe film growth of some oxides as shown, e.g. by Ehrentraut et al. [15] and Romanyuk et al.[17]. Basically, in LPE, a solution consisting of solvent and solute is used to deposit a filmon a single-crystalline substrate so as to preserve the crystallographic information resultingfrom the substrate. Consequently, a ZnO substrate would be the ideal choice to fabricateZnO films with zero lattice and thermal misfits.

LPE from chloride solution is basically water-free and, consequently, hydrogen plays no roleas electrically-active shallow donor (18e35 meV [121]) impurity in ZnO [120].

4.2.1. TechnologyDifferent LPE technologies have been developed for a broad spectrum of semiconductors

and oxides. We applied a dipping technique [18] as described by Levinstein et al. [122]. Theschematic setup is shown in Fig. 26. Experiments were carried out using industrial scale

Fig. 26. Schematic of the LPE setup as employed for ZnO film growth.


LPE furnaces with three heating zones with automatic temperature control giving a temperaturedeviation of �1 K. A lift and rotation unit is mounted above the furnace. The liquid solutioninside the crucible is separated into growth and saturation zones above and beneath the baffle,respectively. We used LiCl (melting Point Tm¼ 605 �C) and a eutectic mixture of 35 mol%NaCle65 mol% CsCl (Tm¼ 486 �C) as the low-temperature solvent. The highest solubilityof ZnO at 650 �C was obtained in LiCl, 9.14� 10�2 g with respect to 1 mol LiCl. The solubil-ity of ZnO in NaCleCsCl is about 5.3� 10�3 g mol�1. Due to the comparatively high solubil-ity in LiCl, we have chosen LiCl as solvent for ZnO for all LPE films grown and discussed inthe remainder of the paper.

A K2CO3 pellet on the bottom inside the crucible served as an oxygen source to form ZnOwith zinc from the zinc source ZnCl2, which was thoroughly mixed with the solute. The follow-ing reaction appears in the crucible:

SolventþZnCl2 þK2CO3/SolventþZnOþ 2KClþCO2[: ð13Þ

This reaction results in the formation of KCl that forms a solid solution with all of the ap-plied solvents. The melting point of the solution is slightly lower for LiCl, NaCl, and CsClwhen used as solvents. A large quantity of the resulting ZnO occurs as crystallites 1e10 mmin size along the c-axis and is deposited at the bottom of the crucible. This was used to controlthe supersaturation of the solution at the beginning of each experiment.

The choice of K2CO3 as an oxygen source has been made due to its easy availability andhigh purity. Fig. 27 compares the results from powder XRD measurements of the deposit asthe product of the reaction of ZnCl2 with Na2CO3, K2CO3, and Na2O. The latter was availableonly in 97% purity. Nevertheless the formation of XRD phase-pure ZnO was confirmed.

All chemicals in our experiments have been used in purities� 99.99%. We used crucibles ofhigh-density alumina and zirconia ceramic. The alumina crucibles never showed any corrosion,i.e. neither a weight change nor coloration. Zirconia, by contrast, is slowly attacked by theZnCl2/LiCl solution. This sometimes caused cracking of the crucible after multiple uses.

Fig. 27. Powder XRD of deposited ZnO microcrystals grown from different oxygen sources: (a) Li2CO3, (b) Na2CO3,

(c) K2CO3, (d) Li2O and (e) Na2O.


52 (2006) 280e335

The substrate was attached to a platinum wire (0.5 mm diameter), which was mounted on analumina rod. The substrate was immersed into the liquid solution after a dwell time of 2 h andthe rotation speed was slowly increased up to 5e10 rpm. The vertical temperature gradient inthe solution between the bottom and the substrate was set to zero.

The experiments were performed under air at atmospheric pressure and at a constant tem-perature. The growth was finished by separating the substrate from the solution. The adhesivesolvent was removed from the film by rinsing with distilled water. After that 2-propanol wasused to remove water from the sample.

The state of the oxygen source, i.e. as a powder of micrometer particle size or as a polycrys-talline pellet, strongly affects the duration of the above chemical reaction [123]. The powderreacted within few minutes whereas the pressed and sintered pellet serves well over 24 h.This can be explained in terms of surface size of the oxygen source, which is very large forthe powder and minimized in case of the pellet.

The alkaline-metal chlorides possess some advantageous properties, which are (a) low melt-ing points from 605 �C for LiCl to 801 �C for NaCl; (b) complete dissociation to monatomicions when molten [124], which results in a low dynamic viscosity; (c) a high solubility in water,therefore cleaning of the grown film is quite easy; (d) availability in high purity at reasonableprice.

The ambiguous growth nature of ZnO along the h0001i direction requires different growthconditions for either the face (0001) or the ð0001Þ. Films grown on ð0001Þ generally were ofpoor quality and much lower thickness. In good agreement to Suscavage et al. [48] we observedthe ratio of growth speed of ð0001Þ : ð0001Þ ¼ ð2:8e3Þ : 1. That is the reason why we focusedexclusively on the growth on (0001) ZnO. In addition, the (0001) face is considered to be su-perior to ð0001Þ with respect to p-type doping [125].

4.2.2. Undoped ZnO filmThe concentration of ZnO determines the growth mechanism [18], which evolves from

island growth at cZnOsol < 2 mmol ZnO to columnar growth (free-standing columns

at cZnOsol ¼ 2.5� 0.5 mmol ZnO, to porous columnar film at 3.5� 0.5 mmol ZnO, to closed co-

lumnar film at cZnOsol ¼ 5 mmol ZnO) thence to step-flow growth at cZnO

sol ¼ 12� 0.5 mmol ZnO.The SEM images (Fig. 28) show: (a) the film grown from 10 mmol ZnO displays a columnargrowth mechanism; (b) 13 mmol ZnO results in large steps, which are aligned parallel toeach other; (c) step bunching is obtained for 16.5 mmol ZnO in the solution.

High structural quality, single-crystalline films, X-ray full-width half-maximum (FWHM) of26e31 arcsec for (0002) reflection, with steps propagating over macroscopic dimensions werefabricated in the concentration range 13� cZnO

sol � 15 mmol ZnO. The X-ray reciprocal spacemap using the (101) reflection (Fig. 29) shows the high crystallinity of the film (b) grownfrom 13 mmol ZnO concentration (a). The FWHMs for 2q/u and u are 22 arcsec for bothand 12 and 15 arcsec for substrate and film, respectively.

The morphological conditions of the substrate are apparently conserved [18]. The stepsthemselves possess a very flat surface with no contrast being resolved by AFM (Fig. 30,scan area 5� 5 mm2). The RMS roughness of 0.218 nm was calculated. The line scan revealsan interstep distance of 2� 0.5 mm and a step height of 0.5 nm, which corresponds to the heightof one unit cell. The thickness of the film is about 1 mm as estimated from weight difference ofthe substrate before and after growth. Typically, the film thickness varies with growth time. Thegrowth rate for film growth on (0001) ranges between 0.12 mm h�1� Vh0001i � 0.25 mm h�1.Best qualities were grown at the lowest value of Vh0001i.


Fig. 28. SEM images of ZnO films grown under different cZnOsol . (a) Columnar growth mechanism is observed for the film

grown from cZnOsol ¼ 1 mol%. (b) The film grown from cZnO

sol ¼ 1.3 mol% displays large steps. (c) The film grown from

cZnOsol ¼ 1.65 mol% is showing the effect of high supersaturation.

Fig. 29. X-ray reciprocal space maps using (101) reflection from (a) the substrate and (b) the homoepitaxial ZnO film

grown on this particular substrate.


52 (2006) 280e335

The films do appear completely transparent, colorless, and highly reflective. Step distanceand height strongly vary with cZnO

sol .Fig. 31 shows the SIMS profile from a Ga-doped, 1.7 mm thick ZnO film on a hydrothermal

ZnO substrate. The normalized intensities of the species Li, Ga, Mg, Al, Na, and Si are shown.The Na and Li counts are increased in the film up to two orders of magnitude in comparisonwith the ZnO host. Typical for the proposed LPE growth of ZnO, cations from the solventwould be detected in the ZnO film by SIMS. The intentionally undoped ZnO films are thereforeslightly doped with alkaline-metal ions (order of 100 ppm wt for Li). Mainly the startingchlorides are made responsible for that, whereas the source for Al and Si might rather bethe ceramic crucibles [18].

4.2.3. Solid solutions and doping

4.2.3.1. MgOeZnO. Magnesium is widely used for band gap engineering of ZnO to achievea UV shift. This shift increases with increasing x in MgxZn1�xO. However, MgO crystallizesin the rock-salt structure (coordination number of nearest neighbors is 6, by contrast to 4 incase of wurtzite structure), which is likely the reason for the increasing tendency for phase seg-regation into MgO and ZnO with increasing x in MgxZn1�xO. Phase segregation for x� 0.36was reported for the growth of thermodynamically metastable films by PLD by Ohtomo

Fig. 30. AFM image and line scan revealing monatomic steps on the surface of an LPE-grown undoped ZnO film.


et al. [117]. The thermodynamic solubility limit of x z 0.04 at 1600 �C [126] has been ac-cepted for a long time. Recently, however, x z 0.06 was reported by our group for films grownby LPE on a (0001) ZnO substrate [127]. The source for Mg was MgCl2 (99.9% purity) whichwas employed at cMg

sol � 20 mol%. The growth occurred at 640 �C using a growth time� 24 h.A micrograph of a film taken by NDIM is shown as Fig. 34a. Growth steps and the formation ofa 1011g�

facet at the rim of the (0001) substrate is characteristic [127].Based on structural analysis using 2q/ueu X-ray reciprocal space maps of the symmetric

(002) and asymmetric ð105Þ reflection from grown MgxZn1�xO/ZnO heterostructures andX-ray powder diffraction from solid phases deposited at the bottom of the crucible, the solubilitylimit has now been verified between 0.06� x� 0.1 [128]. The in-plane strain in films increaseswith increasing x as indicated by broadening of the u peak to 176 arcsec for x¼ 0.06 (Fig. 32a).Using the asymmetric ð105Þ reflection (Fig. 32b) the MgxZn1�xO film grown fromcMg

sol ¼ 10 mol% is characterized by strain relief as the peak from the film is slightly left-hand shifted with respect to the peak from the ZnO substrate. Recently, Wang et al. [50]have reported the hydrothermal growth of ZnO bulk crystals containing up to 5.5 mol% Mg(see also Table 3). This result strongly confirms our findings. In fact, the similarity of both re-sults is striking despite the different chemistry of the solutions.

SIMS measurements from the region near the interface of MgxZn1�xO/ZnO heterostructureshave been carried out and do reveal a highly uniform Mg content within the doped films(Fig. 33). The transition of the Mg concentration between substrate and film is relatively sharp[128]. More growth results are reported in Table 6.

Hall measurement on the Mg0.06Zn0.94O film made under the same conditions as for the hy-drothermal ZnO (Section 3.2.3) revealed a mobility of mH¼ 20 cm2 V�1 s�1 at 560 K andN z 1017. Lorenz et al. [129] reported mH about 20 cm2 V�1 s�1 at 300 K and N z 1017 cm�3

for PLD films grown on sapphire substrates and Ogata [130] gave mH¼ 100 cm2 V�1 s�1 at300 K and N¼mid 1017 cm�3 for MBE grown films on a-plane sapphire.

4.2.3.2. CdOeZnO. It has been shown that the system CdOeZnO can be employed to achievea red shift in the excitonic emission from CdxZn1�xO films grown by PLD, MBE, MOVPE, and

Fig. 31. SIMS profile from a Ga-doped, 1.7 mm thick homoepitaxial ZnO film. The hatched area comprises the cross-

section of the LPE film.

52 (2006) 280e335


52 (2006) 280e335

RPE-MOCVD [8,131e135]. Concentrations of x� 0.697 in the films have been reported,which is far beyond the thermodynamic solubility limit of x¼ 0.02 as reported by Makinoet al. [8]. The latter value, however, is likely the limit for ZnO films fabricated by solvothermalmethods. This might be supported by the fact that the effective ionic radius is considerablylarger for Cd2þ, rIV

Cd2þ ¼ 0:78 A than for Zn2þ, rIVZn2þ ¼ 0:60 A [136]. Moreover, as for MgO,

CdO crystallizes in the rock-salt structure. Consequently, the solubility range is expected tobe narrower than for the system MgOeZnO.

We used anhydrous CdCl2 (99.998% purity) as a Cd source. A typical LPE-grown Cd-dopedZnO film grown from cCd

sol¼ 1 mol% is shown in Fig. 34b. Defects of unknown nature arevisible as dark spots. Films 1.1 and 3.1 mm thick have been grown at Vh0001iz 200 nm h�1

(Table 6). The growth speed under similar growth conditions was about a factor of two higherthan for the other dopants except for Cu.

SIMS measurements do not reveal any significant change of the Cd content betweensubstrate and film. We speculate that the formation of CdO is favored over ZnO. In this caseCdO would precipitate from the solution.

Fig. 32. 2q/u over u X-ray reciprocal space maps of MgxZn1�xO/ZnO heterostructures grown from solutions with dif-

ferent cMgsol using (a) the symmetric (002) reflection. The broadening of the u peak indicates increasing strain in the film

with increasing x. (b) Asymmetric (�105) reflection: the MgxZn1�xO film grown from cMgsol ¼ 10 mol% is characterized

by strain relieve as the peak from the film is slightly left-hand shifted with respect to the peak from the ZnO substrate.


4.2.3.3. Ga2O3eZnO. The Ga dopant in ZnO is very effective in enhancing the electricalconductivity [137]. The effective ionic radius of Ga3þ is smaller than for Zn2þ, rIV

Ga3þ ¼ 0:47A [136].

We used anhydrous GaCl3 (>99.999% purity) as the Ga source. The films appeared withstrongly pronounced growth steps on the (0001) face and large, parallel aligned steps are typical(Fig. 34c). The film was grown from cGa

sol¼ 5 mol% and was 1.7 mm thick. The XRC FWHM of28 arcsec was measured (Table 6).

In contrast to all the other doping attempts, the formation of steps of a few nm in height hasbeen observed for the growth on the ð0001Þ face. This indicates a change of the solubility ofGa-doped ZnO in LiCl in comparison to the system ZnOeLiCl.

1.E+16

1.E+17

1.E+18

1.E+19

1.E+20

1.E+21

1.E+22

0 1000 2000 3000 4000 5000Depth [nm]

undoped1.0 mol%3.0 mol%

Mg

conc

entra

tion

[atm

/cm

3 ]

Fig. 33. SIMS profile from the region near interface of Mg-doped films.

Table 6

LPE growth of doped ZnO films on the (0001) face of hydrothermal ZnO substrate

System csol

(mol%)

cfilm

(mol%)

XRC FWHM

(arcsec)

Thickness

(nm)

Vh0001i(nm h�1)

Remarks

MgO �30 �6 15e176 300e1700 17e109 Increased tendency for facet formation

CdO 0.1 41 1100 214 Some inclusions; growth about two

times faster1 3100 192

Ga2O3 5 28 1700 109 Some inclusions; formation of growth

steps on ð0001Þ face

In2O3 0.2 24 530 33 Extremely low growth speed

2 35 800 50

GeO2 0.525 22 1500 92 Pronounced faceting exhibits 1010g�

faces3.3 31 1900 117

Sb2O5 5 1500 94

CuO 5 17.3 5200 325 Fast growth


52 (2006) 280e335

The concentration of Ga in the ZnO films has not been estimated quantitatively, the SIMSprofile, however, shows an increase in the counts of about two orders of magnitudes for the filmwith respect to the substrate (Fig. 31).

4.2.3.4. In2O3eZnO. Indium in ZnO acts similarly to Ga. Doping ZnO with In led to an in-crease in the n-type conductivity of the ZnO films, which may be used in applications, i.e.as a transparent conductor. The use of In-doped ZnO as a scintillator [138] has recently becomeattractive, particularly in the light of reduced self-absorption effects [97,139]. The effectiveionic radius of In3þ is slightly larger than for Zn2þ, rIV

In3þ ¼ 0:62 A [136].InCl3 (99.999% purity) has been employed for use as the In source. The XRC FWHM using

the (0002) reflection (Fig. 35b) gives 24 and 35 arcsec for the films of thicknesses 530 and800 nm grown from cIn

sol¼ 0.1 and 2 mol%, respectively (Table 6). The growth speed is morethan a factor of two lower than for the comparable Ga-doped film.

4.2.3.5. GeO2eZnO. Only a few reports attempted to change the electronic structure of ZnO bydoping with Ge [140]. Ge-doped ZnO ceramic was obtained by solid-state reaction and somefilms have been grown by PLD by Fan et al. [141]. The occupation of the Zn site shoulddecrease the lattice parameters (rIV

Ge4þ ¼ 0:39 A [136]) while it would lead to an increasewhen present as an interstitial. Both cases, however, yield lattice distortion. The solubility ofGe in ZnO is rather limited to about 0.7 mol% [140]. We tried to incorporate Ge4þ intoZnO by LPE.

Fig. 34. NDIM micrographs taken from films grown with different dopant concentration in the initial solution:

(a) 6 mol% Mg, (b) 1 mol% Cd, (c) 5 mol% Ga, (d) 3.3 mol% Ge.


The source for Ge4þ was GeI4 (99.999% purity). We worked with cGesol¼ 0.525 and 3.3 mol%

at constant growth temperature of 640 �C. Fig. 34d shows the NDIM image of the film grownfrom cGe

sol¼ 3.3 mol%, which reveals a pronounced formation of 1010g�

facets. Vh0001i wasbetween 92 and 117 nm h�1 (Table 6). The FWHM increased from 22 arcsec for0.525 mol% doping to 31 arcsec for 3.3 mol% doping (Fig. 35a).

4.2.3.6. Sb2O5eZnO. Antimony (Sb) might be a candidate as p-type dopant in ZnO if oxygensubstitutes for Sb on the oxygen site. The effective ionic radius of Sb5þ is rVI

Sb5þ ¼ 0:62 A and ofO2� rIV

O2� ¼ 1:38 A [136].The source for Sb was ZnSb (99.999% purity). We worked with cSb

sol¼ 5 mol% at a constantgrowth temperature of 640 �C. A film of 1.5 mm has been grown on the (0001) face,Vh0001i ¼ 94 nm h�1 (Table 6). The surface appears rather rough in comparison to the

Fig. 35. XRC measurement reveals the broadening of the (0002) reflection peaks as result of increased doping concen-

tration from (a) 0.5 to 3 mol% Ge; film thickness 1.5 and 1.9 mm, respectively. (b) 0.1 and 2 mol% In was doped; film

thickness 0.53 and 0.80 mm, respectively.


52 (2006) 280e335

Mg0.06Zn0.94O film of Fig. 34a. The XRC FWHM from the (0002) reflection was about 30 arc-sec. The sign of a shoulder formation at the lower angle has been found, which might be due tothe phase separation.

The SIMS measurement implies significant incorporation of Sb as the count intensity in-creases by a factor of about 103. The interface between the substrate and the film was hardlydetectable as the count intensity increased slowly.

4.2.3.7. CuOeZnO. The source for Cu was anhydrous CuCl2 of (99.99% purity). We workedwith cCu

sol¼ 5 mol% at constant growth temperature of 640 �C. The growth rate Vh0001i was325 nm h�1 (Table 6). The XRC measurement using the (0002) reflection gives an FWHMof 17.3 arcsec. This was the lowest value for an FWHM we have ever measured from anLPE-grown ZnO-based even though the doped film was grown from a higher supersaturation,cZnO

sol ¼ 17.5 mmol.SIMS revealed that Cu had been easily incorporated into the ZnO film. The count intensity

raised two orders of magnitude and a clear interface between substrate and film was observed.This led us to conclude that Cu-doping of LPE ZnO films is easy to achieve.

4.2.4. Photo- and radioluminescencePL was measured on undoped and doped ZnO LPE films at 4 K under similar conditions as

for the bulk ZnO, see Section 3.2.3. The signal from an undoped ZnO film grown on thehydrothermal ZnO substrate from TEW is shown in Fig. 36a. The dominant peak signal liesat 3.3601 eV with a shoulder peak at 3.364 eV [18]. This is attributed to neutral donor boundexcitation (D0X). More emission peaks at 3.372, 3.374, 3.378, 3.39, and 3.42 eV have beenmeasured. Among them, 3.378 eV is assigned to the A-free exciton (FEA) from the groundstate, 3.39 eV due to B-free exciton (FEB), and 3.42 eV due to n¼ 2 state (or 1st excited state)of FEA [88,125]. The latter has not been measured from the substrate. It was reported [142] thatFEA emission structure becomes clearly visible at higher TG, i.e. 630 �C. Apparently, our resultis comparable to the ZnO films grown on ScAlMgO4 [142], which appear with similar splittingof exciton ground states. From reflectance measurements we determined the binding energiesfor FEA and FEB of 58 and 55 meV, respectively.

The substitution of Zn for Mg in the ZnO lattice shifts the excitonic emission to higher en-ergies [143] as shown in Fig. 36b. The two peaks at 3.357 and 3.425 eV are related to the ZnOsubstrate and the Mg0.06Zn0.94O film, respectively [127]. The inset in Fig. 36b shows the PLemission from 1.8 to 3.7 eV. The broad band peaks around 2.4 eV. The peak from the substratelies at 3.359 eV before being used for the epitaxial growth, see the dashed line. It has beenshown by Meyer et al. [5] that the intensity of the emission peak around 3.36 eV decreasesupon annealing at 600 �C. By contrast, the bound exciton line at 3.357 eV remains unaffected.Wang et al. [50] found the excitonic peak at 3.414 and 3.4468 eV for their MgxZn1�xO hydro-thermal crystals with x¼ 0.03 and 0.055, respectively.

The Cd-doped film (cCdsol¼ 1 mol%) exhibits PL emission peaks only from the ð0001Þ face at

3.36, 3.346, 3.295, 3.225, and around 3 eV. The latter is possibly due to the slightly increasedCd content. It was shown that a Cd0.04Zn0.96O film grown by MOCVD shows a characteristicpeak at 2.91 eV [132]. We assume that the Cd concentration in the (0001) film is below thedetection limit but is well above it in the ð0001Þ film of lower crystallinity.

The situation is different when the film was doped with Ga as shown in Fig. 37a. The boundexciton emission spectrum is considerably changed and the dominant peak becomes the I8 at3.358 eV, which is not observed from undoped ZnO. A slight red shift by about 0.5 meV was


explained by electroneelectron interactions [5] causing an increased metallic conduction behav-ior. Also in good agreement with findings from other groups [5,143] is the greatly weakenedemission at 3.369 eV. The TES peak is at 3.32 eV, which is in good accordance with Ref. [5].

The effect of In-doping is shown in the PL measurement taken at 12 K (Fig. 37b). Thedominant line is the recombination I9 [5] at 3.359 eV and the TES peaks around 3.309 eV.

Not shown here are the PL measurements from LPE films doped with Ge, Sb and Curecorded at 12 K. In the Ge-doped film the major peak was around 3.357 eV, similar to thatin the hydrothermal ZnO substrate which was used in this experiment. The emission peak at3.372 eV was weakening with increasing cGe

sol from 0.525 to 3.3 mol%. The reason remainsunknown at present.

Fig. 36. Low-temperature PL spectra from the region near the band edge of (a) the undoped ZnO substrate (dashed line)

and the LPE-grown film (straight line) and (b) the undoped ZnO substrate (dashed line) and an MgxZn1�xO film

(straight line). The peaks at 3.357 eV and 3.425 eV are due to emission from the ZnO substrate and the MgxZn1�xO

film, respectively. The inset shows the PL spectrum of the MgxZn1�xO film over the energy range from 1.8 to

3.7 eV. Intensity is Log scale for all graphs.


52 (2006) 280e335

Sb-doping reveals peaks at 3.357 and 3.372 eV and the DAP peaks at 3.31 eV. At lowerenergies there is a peak at 3.265 eV with replicas repeating at 72e75 meV, i.e. at 3.193,3.12 and 3.045 eV.

The Cu-doped film showed the peaks at 3.371, 3.358, 3.31 and 3.235 eV.RL measurements at temperatures from 80 K up to 295 K have been obtained following

X-ray exposure at lexc¼ 0.154 nm, using a spectrofluorometer (199S, Edinburgh Instruments)equipped with a steady-state X-ray excitation source (35 kV) and combined with a single grat-ing monochromator and with the photomultiplier XP2233 used in the photon counting mode.Fig. 38a details the spectrum from the bulk ZnO (dotted line) and a typical LPE-grown film(straight line) of 2 mm in thickness. Luminescence from the band edge region is almost sup-pressed in the bulk sample and the ratio of the arbitrary intensity to the broad band emissionfrom the bulk sample peaking around 2.1 eV is only about 0.1. The LPE film shows contrarystrong emission with a maximum around 3.2 eV. The ratio of intensity from this emission to

Fig. 37. Low-temperature PL spectra from the region near the band edge of (a) a 5% Ga-doped and (b) a 0.1% In-doped

ZnO film. Intensity is Log scale for all graphs.


Fig. 38. Room-temperature RL spectrum of (a) an undoped LPE ZnO film and a hydrothermal grown bulk sample of

ZnO. The integral under the curve between 3.1 eV and 3.18 eV contributed to the measurement of the decay time (b,c).

Room-temperature RL decays from (b) a undoped ZnO and (c) an In-doped ZnO sample. Intensity is Log scale for

graphs (b) and (c).


52 (2006) 280e335

the broad band peak around 2.3 eV behaves like 2:1. The observed effect is attributed to a dam-aged surface layer due to machining [18,97] and, by contrast, is not observed by XRC wheresubstrate and film show comparable FWHM.

ZnO holds great expectations as scintillator material, particularly its very short luminescencedecay time t in the sub-nanosecond range is of interest for applications in time-resolved de-vices like time-of-flight positron emission tomography (TOF-PET). Such TOF-PET devicesin medical scanning applications could for example enable the detection of the initial stateof a cancer due to the high spatial resolution available. Up to now, barium fluoride (BaF2) isthe hottest candidate as a super-fast TOF-PET device. Very fast luminescent decay of BaF2

of 0.6 ns was reported [144]. Unfortunately, the density of BaF2 (4.88 g cm�3) is rather lowwhich may hinder its widespread application. High-density material, however, is required toprovide high radiation stopping power so that high detection efficiency can be achieved.Consequently, the thickness of the detector can be lowered by using a high-density scintillatormaterial, which is now a strong requirement from industry. Moreover, while BaF2 emits around200 nm, the exciton wavelength of ZnO is around 400 nm which can be easily detected bya conventional photomultiplier. For Ga-doped ZnO powder the FWHM of the output pulseof 0.21 ns at 295 K was reported by Derenzo et al. [145] and about 0.65 ns FWHM upon ex-posure to a 241Am radiation source was demonstrated for an In-doped ZnO melt-grown crystalby Simpson et al. [138]. Doping with the group-V metals gallium and indium serves to decreaseself-absorption at higher energies beyond 3.36 eV [97].

Luminescence decay time was measured for the energy region 3.1e3.18 eV, i.e.400� lem� 390 nm, which refers to the integral intensity as marked by the shaded area inFig. 38a. An undoped, epi-ready ZnO wafer and the In-doped as grown platelet were exposedto a pulsed laser beam (160 fs pulse duration, lexc¼ 260 nm, 1 kHz repetition rate, excitationenergy< 1 mJ/pulse, 109 W cm�2 peak intensity). The decay spectrum of the undoped and theIn-doped sample is shown in Fig. 38b and c, respectively. Two decay components have beenassigned for both samples: a fast one with 23 and 40 ps decay time and a slow one with 100and 647 ps decay time for undoped and In-doped ZnO, respectively. The measured curvesare fitted by a second exponential decay with offset: I(t)¼ 2.2 e[�t/0.023 ns]þ 0.057e[�t/0.1 ns]þ 0.00024 in the case of an undoped ZnO wafer and I(t)¼ 0.315e[�t/0.04ns]þ 0.92 e[�t/0.647ns]þ 5 e�5 for the In-doped ZnO platelet. The rather noisy spectrumobtained from the undoped, epi-ready ZnO wafer for t> 0.2 ns is speculated to be due toa damaged surface layer as already discussed before. Consequently, emission quenching islikely, which causes the very weak signal and non-radiative energy transfer is strongly pro-nounced [97].

The better understanding of the scintillator properties of ZnO and the effects of doping andmachining have been emphasized in our group since recently and more output may be expectedin the forthcoming years. LPE is being used in this context as a tool of fast screening to obtainZnO-based, thermodynamically stable phases with surfaces free of any mechanical damage dueto machining. Also, the role of DAP recombination in the luminescence from LPE-grown ZnOfilms is now being investigated.

5. Summary and future trends

Solvothermal technologies for ZnO crystals and films were reviewed. Among all the solventsinvestigated supercritical water turned out to be superior in terms of crystal perfection, processcontrol and scalability. The hydrothermal technique produces the largest ZnO single crystals


with excellent structural properties. Two-inch size is now available from mass production and 3inch size has already been demonstrated in this paper. It would be expected that the trend toincrease ZnO crystal size and output will accelerate with increasing demand from industry.This would certainly contribute to a significantly lower price for ZnO wafers. However, criticalissues remain like the quality of the surface of wafers. Manipulation of effects connected withsurface conductivity is necessary. Doping of hydrothermal ZnO could open more applicationwindows such as its role as a scintillator. In particular reproducible p-type doping is expectedto boost ZnO technology.

The fabrication of ZnO films from aqueous and chloride solutions may gain competitivenessto the growth from vapor phase for special applications in ZnO thin film technology. This isdemonstrated for the case of homoepitaxy by LPE, where highly-crystalline ZnO films canbe fabricated from alkaline-metal chloride solutions at temperatures <650 �C. A hydrothermaland alkaline-metal free route produces thick ZnO films on foreign substrates as low as 90 �C.Moreover, solution-grown ZnO films have the potential for in-situ doping since the high crys-tallinity of the film is preserved. Still, there are many open questions asking for continuation inboth fields, the growth of large bulk crystals and films from liquid solutions.

Acknowledgements

The authors are grateful to T. Ono and K. Maeda (Tokyo Denpa Co. Ltd., Tokyo) for pro-vision of ZnO substrates, E. Ohshima and J.M. Ko (both formerly with Tohoku Univ.) for ex-periments, M. Miyamoto (Mitsubishi Gas, Tokyo) for some LPE experiments and SIMSmeasurements, B.-P. Zhang and N.T. Binh (Photodynamics Research Center/RIKEN, Sendai)for assistance in PL measurements, M. Nikl (Institute of Physics, AS CR, Prague) for RL in-vestigation, P. Kiesel and O. Schmidt (Palo Alto Research Center Inc., Palo Alto) for Hall mea-surements and discussion, K. Sugiyama (Univ. of Tokyo) for measuring X-ray topography andT. Yao (Tohoku Univ.) for using the AFM. D.E. expresses his appreciation to F.F. Lange (Univ.of California, Santa Barbara) for discussion and providing data and thanks F. Orito (MitsubishiChemical Corp., Tokyo) for support. Finally, the great help of J.B. Mullin in the final revision ofthe manuscript is deeply appreciated.

We gratefully acknowledge funding by following organizations: The Japanese IndustrialTechnology Research Grant Program in 03A26014a from New Energy and Industrial Develop-ment Organization (NEDO); The Special Coordination Fund by the Ministry of Education,Culture, Sports, Science; The technology program ‘‘Development of Growth Method ofSemiconductor Crystals for Next Generation Solid-State Lighting’’.

Particular gratitude is due to Mitsubishi Chemical Corp. and Tokyo Denpa Co. Ltd. forsupport over the years.

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Dr. Dirk Ehrentraut was born in Germany in 1968. He was educated at Humboldt University

of Berlin, Germany where he obtained a Diploma in crystallography in 1995. From 1995 to

1997 he was Researcher at the Institute of Crystal Growth, Berlin. In 1997 he joined the In-

stitute of Micro- and Optoelectronics at the Swiss Federal Institute of Technology Lausanne

(EPFL), Switzerland. From 2000 to 2003 he worked at the Institute of Applied Optics of

the EPFL from where he obtained a Doctorate in Science for his thesis on novel coherent light

sources. Since end of 2003 he is invited as Visiting Associated Professor at the Institute of

Multidisciplinary Research for Advanced Materials (IMRAM), Tohoku University, Sendai,

Japan, where his focus is on crystal chemistry and growth of wide band gap materials

(ZnO, III-Nitrides) and sesquioxides fabricated from liquid solution. His field of competence

includes the growth of bulk crystals and single-crystalline films of a large variety of oxides and

semiconductors applying vapor phase and liquid phase (melt, solution) technologies.

Hideto Sato was born in Japan in 1973. He graduated from the Ritsumeikan University in

Shiga, Japan and afterwards joined Murata Manufacturing Co., Ltd. where his research was

focused on piezoelectric bulk crystals. From 2004 on he worked on the growth of ZnO film

within the frame of a joined project between Murata Manufacturing Co., Ltd. and IMRAM.

Yuji Kagamitani was born in 1977 in Japan. He obtained a Masters in Science (MSc.) in 2002

from the Department of Chemistry, Faculty of Science, Tohoku University, Sendai, Japan. He

is currently a PhD student at IMRAM, Tohoku University, preparing a PhD thesis entitled

‘‘Study on Crystal Growth of Zinc Oxide with Metal Ion Doping and Its Scintillation Proper-

ties’’. His research interest is on the solvothermal growth of ZnO and GaN crystals.

Dr. Hiroki Sato was born in 1974 in Yamagata/Japan. He graduated with a MSc. from the

Department of Metallurgy, Graduate School of Engineering, Tohoku University in 1997 and

obtained a Doctorate from Tohoku University in 2005. Between 1997 and 2003 he was em-

ployed by NEC Tokin Corp. where he worked on CdMnTe/CdMnHgTe solid solution and rutile

for optical applications. In 2003 he joined Fukuda X’tal Lab. Inc. to continue his research on

fluorides and scintillator materials.


52 (2006) 280e335

Prof. Akira Yoshikawa was born in Japan in 1970. He obtained a MSc. and Doctorate in Sci-

ence from the University of Tokyo in 1996 and 1999, respectively. He became Research

Associate of the Fukuda laboratory at the Institute of Materials Research (IMR), Tohoku

University, Sendai in 1997. In 2002 he was assigned Research Associate at IMRAM, Tohoku

University, Sendai. Since October 2003 he is an Associate Professor at IMRAM and temporary

a Visiting Associate Professor of the University Claude Bernard Lyon 1, France.

His research interests are on the melt growth of single-crystalline oxides and fluorides for scin-

tillator and laser applications and the solvothermal growth of wide band gap semiconductors

ZnO and GaN.

Prof. Tsuguo Fukuda was born in Japan in 1939. In 1964 he graduated from the Faculty of

Science, University of Tokyo, Japan. During 24 years of working in Tokyo Shibaura Electric

Co. (named later Toshiba Corp.) including 5 years working in the Optoelectronic Joint Re-

search Lab. in Japan, in 1971 he obtained a PhD from the University of Tokyo for his thesis

on ferroelectric crystals. He became Professor of IMR, Tohoku University in 1987. In 2002,

he became Research Professor of IMRAM, Tohoku University. His research interests are crys-

tal growth technology of bulk single crystals for optical applications. Presently, one of his main

research topic is hydrothermal and ammonothermal growth of ZnO and GaN. He is author or

co-author of more than 600 papers. He is the founder and president of Fukuda X’tal Lab. Inc.

Documents

Crecimiento Solvotermal de ZnO