reaction mechanism between cu zn sn se components for …mechanism... · 2.2 ellingham diagram for...

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Reaction mechanism between

Cu–Zn–Sn–Se components for the

formation of Cu2ZnSnSe4 film

Der Naturwissenschaftlichen Fakultäten

der Friedrich-Alexander-Universität

Erlangen-Nürnberg

zur

Erlangung des Doktorgrades Dr. rer. nat.

vorgelegt von:

Hyesun Yoo

aus Incheon, Republic of Korea

Als Dissertation genehmigt

von der Naturwissenschaftlichen Fakultät

der Friedrich-Alexander-Universität Erlangen-Nürnberg

Tag der mündlichen Prüfung: 23.06.2016

Vorsitzender des Promotionsorgans: Prof. Dr. Jörn Wilms

Gutachter/in: Prof. Dr. M. Alexander Schneider

Gutachter/in: Prof. Dr. Rainer Hock

Gutachter/in: Prof. Dr. Michael Thoss

This study is supported by SGR in Paris and AVNCIS in München.

Zusammenfassung

Das Herstellen von homogenen kesterit-basierten Cu2ZnSnSe4 Dünnschichten

erfordert ein tiefgehendes Verständnis der Reaktionsmechanismen zwischen vier

Elementen. Zu diesem Zweck werden verschiedene Proben mithilfe der zeitaufgelösten

Röntgenpulverdiffraktometrie untersucht, und so die Reaktionspfade der einzelnen

Elemente ermittelt. Um die einzelnen Reaktionen voneinander zu trennen werden Proben

mit unterschiedlicher Anzahl metallischer Schichten (Ein-, Zwei- und Dreimetallproben)

hergestellt. Die gemessenen Diffraktogramme werden basierend auf den entsprechenden

Phasendiagrammen der Legierungen und den Referenzdiffraktogrammen der einzelnen

Phasen ausgewertet. In einigen Fällen kommen zusätzlich Röntgenpulverbeugung unter

streifendem Einfall und Ramanspektroskopie zum Einsatz.

Die Reaktionen der Einmetallproben zeigen den Temperaturbereich, in dem ein

bestimmtes Metall eine Verbindung mit Se eingeht, ohne den störenden Einfluss anderer

Elemente. Basierend auf diesen Ergebnissen lassen sich die Reaktionstemperaturen aller

binären Selenide, die in den Zweimetallproben entstehen, miteinander vergleichen. Dieser

Vergleich kann den Einfluss anderer Elemente auf die Bildung der binären Selenide

aufzeigen.

Die Ergebnisse der Zweimetallproben lassen vermuten, dass verschiedene

Bildungsmechanismen eine Rolle spielen, abhängig von der Stapelfolge im Präkursor,

auch bei nur zwei metallischen Schichten. Folglich hängt die Reihenfolge der

Legierungsreaktionen stark von der Stapelfolge ab. Aus der Analyse der verschiedenen

Reaktionspfade lassen sich Rückschlüsse auf die Reaktionspräferenzen der einzelnen

Elemente und Bildungsmechanismen der Legierungen ziehen.

Mithilfe dieser Erkenntnisse können die Reaktionspfade der Dreimetallproben und die

Bildungsmechanismen von CZTSe untersucht werden. Entsprechend sind die

Dreimetallproben für die genauere Analyse je nach Bildungsmechanismus in

Unterkategorien eingeteilt.

Die Bildungsmechanismen für jeden Bestandteil der CZTSe Synthese sind in Kapitel

4.3 zusammen mit den Restphasen, die in der Schicht zurückbleiben, einzeln aufgeführt.

Durch die in dieser Arbeit herausgearbeiteten Reaktionscharakteristika lässt sich eine

optimale Stapelfolge für den vorgeschlagenen Modellpräkursor vorhersagen.

Summary

To find an approach to prepare a homogenous kesterite-based Cu2ZnSnSe4 thin film, it

is necessary to understand the reaction mechanism between four elements. Thus, various

samples are analysed by time-resolved in situ X-ray powder diffraction (XRD) to observe

the reaction process of each element. To better understand the formation reactions,

samples with different number of metallic layers (one-, two- and three-metal samples) are

prepared. The obtained diffractograms are analysed on the basis of the relevant alloy phase

diagrams and XRD reference patterns of each phase. When necessary, grazing-incidence

XRD measurement and Raman spectroscopy are performed on several samples.

Reaction paths of one-metal samples show the formation temperature of a metal with

Se without any disturbance by other elements. On the basis of the results, the formation

temperature of each binary selenide in the results of two-metal samples will be compared.

This comparison may show the influence of other elements on the formation of binary

selenides.

Results for two-metal samples suggest different formation processes depending on the

sequence of stacking layers in the initial precursor, although these samples include only

two metallic layers. This means that the stacking order of precursors significantly affects

the sequence of alloy formation. From the different reaction paths, the tendency of four

elements to react may be determined together with several characteristics of formation

reactions of the alloy phases.

Based on these analyses, the reaction paths of three-metal samples may be revealed,

and several characteristics of formation reactions of CZTSe may be observed. Therefore,

the three-metal samples are sub-divided again according to these characteristics for

detailed analysis.

Formation processes of each component for the CZTSe formation are separately

described in section 4.3, along with the cause of the remaining secondary phases in the

kesterite film.

The reaction characteristics revealed in this study provide information on the optimum

stacking order of precursors and lead to a conclusion on one proposed precursor.

Contents

1. MOTIVATION 1

2. MATERIALS AND METHODS 2

2.1 Phase diagram of the alloys 2

2.1.1 Metal system 2

2.1.1.1 Cu–Zn binary alloy 2

2.1.1.2 Cu–Sn binary alloy 3

2.1.1.3 Zn–Sn binary alloy 3

2.1.1.4 Sn–Zn–Cu ternary alloy 3

2.1.2 The selenium–metal phase diagram 4

2.1.2.1 The Cu–Se thermodynamic system 4

2.1.2.2 The Zn–Se thermodynamic system 5

2.1.2.3 The Sn–Se thermodynamic system 5

2.2 Ellingham diagram for Se 5

2.3 Crystal structure data for the observed phases 6

2.4 X-ray powder diffraction analysis 7

3. EXPERIMENTAL DETAILS 11

3.1 Preparation of the Cu2ZnSnSe4 precursor 11

3.2 Characterisation techniques 12

3.2.1 Time-resolved in situ X-ray diffraction 12

3.2.2 Grazing Incidence X-ray Diffraction 16

3.2.3 Raman spectroscopy 17

4. RESULTS AND DISCUSSIONS 18

4.1 Investigation of single and binary metal systems with Se 18

4.1.1 Reaction path of a single metal with Se 18

4.1.1.1 Reactions of Mo/Zn/Se 18

4.1.1.2 Reactions of Mo/Sn/Se 19

4.1.2 Reaction path of two metals with Se 22

4.1.2.1 Reactions of Mo/Cu/Sn with Se 22

4.1.2.2 Reactions of Mo/Sn/Cu with Se 25

4.1.2.3 Reactions of Mo/Zn/Cu with Se 28

4.1.2.4 Reactions of Mo/Cu/Zn with Se 31

4.1.2.5 Reactions of Mo/Sn/Zn with Se 37

4.1.2.6 Reactions of Mo/Zn/Sn with Se 40

4.1.3 Results of experiments on single and double metal layers 42

4.1.3.1 Influence of pressure on Sn–Se alloy formation 43

4.1.3.2 Reactive Cu 43

4.1.3.3 Outward diffusion of Cu: blocked only by a Zn layer 44

4.1.3.4 Induced movement of Zn to the back electrode by Cu 44

4.1.3.5 Reaction sequence of Cu–Se alloys 45

4.1.3.6 High affinity of Se to Cu 46

4.1.3.7 Delayed ZnSe formation by the Cu layer beneath the Zn layer 47

4.1.3.8 Delayed SnSe formation by Cu and Zn contents of alloy 48

4.1.3.9 Crystalline CuxSey phase determines the rate of Cu2SnSe3 formation 49

4.1.3.10 Conclusion: The tendency of four elements to react with each other 49

4.2 Investigation of ternary metal systems with Se 51

4.2.1 Correlation of delayed ZnSe crystallisation with the reaction sequence of selenides 51

4.2.1.1 Reactions of Mo/Cu/Sn/Zn/Se 51

4.2.1.2 Reactions of Mo/Zn/Sn/Se/Cu 56

4.2.1.3 Reactions of Mo/Zn/Sn/Cu/Se 59

4.2.1.4 Detection of residual ZnSe by Raman scattering 66

4.2.1.5 Discussion 67

4.2.2 Different formation process of two samples with reversed elemental stacking order 70

4.2.2.1 Reaction in Mo/Zn/Sn/Cu/Se and Mo/Se/Cu/Sn/Zn with reversed stacking

order 70


4.2.3 The effect of two Cu layers on the reaction 75

4.2.3.1 Reactions of Mo/Zn/Cu/Sn/Se/Cu 76

4.2.3.2 Reactions of Mo/Cu/Zn/Cu/Sn/Se 82


4.3 Secondary phases in Cu2ZnSnSe4 films: How and why they form 90

4.3.1 Cu2Se 90

4.3.2 ZnSe 91

4.3.3 SnSe2 (or SnSe) 93

4.3.4 Cu2SnSe3 94

4.3.5 Cu2ZnSnSe4 95

5. CONCLUSIONS 96

REFERENCES 98

APPENDIX 105

Supplementary information 105

List of publications 106

Conference contributions 108

ACKNOWLEDGEMENTS 110

1. Motivation

1

1. Motivation

Techniques which convert solar energy into electricity are attractive for as long as the

sun shines in the sky. Because the source of electric energy is unlimited and free,

photovoltaics has been widely studied and has been used in commercial application; at

present, one can buy a solar panel mainly produced from Si. Meanwhile, chalcopyrite-

based Cu(In,Ga)(S,Se)2 (CIGS) solar cells have been of interest because chalcopyrite has

a direct band gap, whereas monocrystalline Si has an indirect band gap. Thus, the

thickness of the absorber layer in the photovoltainc (PV) module can be reduced from

hundreds to 1–2 micrometers. With the efficiency of CIGS-based solar cell of 21.7% [1],

this material demonstrates its potential use as efficient thin film in PV modules. However,

because indium and gallium are expensive, there have been attempts to replace these two

components with zinc and tin, which are cheap and abundant materials: the birth of a new

material, Cu2ZnSn(S,Se)4 (CZTSSe).

The first trial synthesis of Cu2ZnSnS4 (CZTS) film without Se was performed in 1988

by Ito and Nakazawa [2]. They found that CZTS has the appropriate properties for the

absorber layer of the PV module. Afterwards, since a low-cost solar cell is required for

commercial applications, the fabrication of CZTS-based solar cell has been steadily

attempted, including selenium (CZTSSe) to optimise the band gap. In 2013, Wang et al.

proved the feasibility of CZTSSe by producing the CZTSSe-based solar cell with an

effieicncy of 12.6% [3]. As this material holds promise for low-cost solar cells, synthesis

of homogeneous CZTSSe films has becomes one of the main approaches to achieve high

efficiency. Because five elements are used to form this material, a growth of CZTSSe

single crystal film is difficult. After finishing the synthesis of kesterite film, secondary

phases such as ZnS/ZnSe or Cu2S/Cu2Se are easily remaining in the film.

For this reason, understanding the reaction mechanism of this material is necessary.

Some of notable discoveries in this regard have been made. Schurr et al. observed the

reaction path of CZTS depending on the metal ratios in as-deposited films [4]. Weber et al.

found the Sn loss from Cu-Zn-Sn-S films [5], inducing the formation of secondary phases

due to the lack of Sn in the film. These results for the properties of the CZTSe formation

observe only with respect to the proportion of its components. On the contrary to this,

other studies demonstrate that the CZTS formation is affected by the precursor’s order of

stacked layers [6, 7]. Although the elemental ratios of components in these studies were

the same, the completed CZTS films were different [6, 7]. Thus there is a need to

understand the tendency for reaction between the four components of CZTS. This thesis

focuses on elucidating the characteristics of the reaction path between metallic elements

and selenium in relation to the sequence of stacked layers in the precursor because the

reaction process of CZTS is already studied in [4].

2. Materials and Methods

2


2.1 Phase diagram of the alloys

The observation of phase diagrams is necessary to beter interpret the reaction path of

alloys in section 4. There are seven phase diagrams for the alloys in the four components

of CZTSe. They show the equilibrium phases in the alloys according to temperature and

the elemental ratios of each compound. Because experiments in this dissertation are at

temperatures up to 550 C, the phase transition of alloys is observed up to around 550 C.

2.1.1 Metal system

2.1.1.1 Cu–Zn binary alloy

In the phase diagram for Cu–Zn [9], the Cu–Zn alloy, which is termed ‘brass’, is

formed within five kinds of phases at room temperature, depending on the concentration

of Zn (or Cu): α, β’, γ, ε and η phases [8, 9]. Standard formulas of β’ and γ phases are

CuZn and Cu5Zn8, respectively. The α, ε and η phases may correspond to many types of

brass. The α phase denotes Cu-rich CuZn, which includes more copper atoms than those

in β’-CuZn; these include Cu2Zn or Cu0.7Zn0.3 [10]. The ε and η phases represent Zn-rich

CuZn, including Cu0.7Zn2 [11] and Cu0.025Zn0.975 [12], respectively, as standard formulas.

A notable property of Cu–Zn alloy is the tendency of all of its phases to undergo

dezincification. This process is selectively leaching of Zn from the Cu–Zn alloy and is

continuously going on with an increasing rate as the temperature rises. Accordingly, most

boundary lines of the solid phases in the Cu–Zn phase diagram are curved, in contrast to

other phase diagrams for Cu–Se or Cu–Sn. At the boundary line of the α phase near ~33 at%

of Zn, the graph tilts to higher concentration of Zn until 38.27 at% of Zn as temperature

rises by 454 C [9]. It denotes that the required concentration of Cu for the formation of a

pure α phase gradually decreases as the temperature increases for 227–454 °C because of

dezincification. In contrast, the Cu-49 at% Zn alloy transforms from a pure β’ phase

(CuZn) into coexisting β’ and γ phases when this brass is heated up to 468 °C, in

accordance with the tilted boundary line of the β’ phase [8, 9]. This signifies that

selectively leached Zn from the CuZn phase adheres onto a part of CuZn again, so that the

part of CuZn transforms into Cu5Zn8, leaving decomposed Cu: 8 CuZn → Cu5Zn8 + 3 Cu

(partial reaction).

2.1 Phase diagram of the alloys

3

2.1.1.2 Cu–Sn binary alloy

Five kinds of Cu–Sn alloys (η’, η, ε, δ and γ phases) in the Cu–Sn system are observed

up to ~550 °C over the full range of Sn (or Cu) concentration in the Cu–Sn phase diagram

[13-15]. Each of these phases needs a certain temperature to compound. Two η’-Cu6Sn5

and ε-Cu3Sn phases can form at room temperature. Especially the η’-Cu6Sn5 phase

converts into η-Cu6Sn5 at 186–189 C from a superstructure to the hexagonal NiAs (hP4)

structure with the same chemical formula [13]. This converted η phase decomposes at

415 C (or 408 C [14]) into the ε-phase (Cu3Sn) and a liquid Sn [13, 15]. The ε phase

(Cu3Sn) does not transform into another structure until 640–676 C [13, 15] (or until 649–

676 °C [14]), unless the Sn (or Cu) concentration changes.

In the Cu-rich region of this diagram [13, 15], the δ phase (Cu41Sn11) and the γ phase are

compounded at 350 C and 520 C, respectively. Here the γ phase has the same formula

as that of ε (Cu3Sn) but has different structure. In the contrast, a distinct phase other than

Cu6Sn5 does not form in the Sn-rich region of the phase diagram, but only the melting of

Sn occurs from 232 C.

2.1.1.3 Zn–Sn binary alloy

Because no Sn–Zn structure naturally exists, pure metallic Zn and Sn separately

coexist in the Sn–Zn mixture until 198.5 C [16]. At 198.5 C, Zn and Sn comprise a

eutectic alloy when the Zn concentration is higher than 14.9 at%. In other words, the

eutectic composition is 85.1 at% Sn and 14.9 at% Zn, and the eutectic temperature is

198.5 C [16].

Upon heating, this alloy becomes a liquid phase at a certain temperature, depending on

the Sn concentration. As the proportion of Sn increases in the Sn–Zn alloy, the liquidus

temperature becomes lower, and vice versa. When the Sn concentration is 31.65–62.97

at%, the range of the liquidus temperature is 356.8–296.8 C [16]. In particular, the

liquidus temperature is 326.8 C when the Sn concentration is 50.54 at% [16]. Therefore,

one can recognise that the proportion of Sn is lower than 50.54 at% if the eutectic alloy

melts at a temperature higher than 326.8 C. For example, if the eutectic Sn–Zn alloy

melts at ~350 C, then the Sn composition at ~37 at% (Sn-63 at%Zn) may be deduced

from this phase diagram. In the same manner, the liquidus temperature at ~300 C

indicates a Sn concentration of ~62 at% (Sn-38 at% Zn).

2.1.1.4 Sn–Zn–Cu ternary alloy

C. Chou and S. Chen studied the phase equilibria of the Sn–Zn–Cu ternary system at


4

210, 230 and 250 C [17]. According to general observations of the compounding phases

depending on the composition of the three elements, the Cu–Sn alloy does not form a

compound unless the CuZn phase forms. Only when the Cu concentration is higher than

~75 at% (eg, a Sn-15 at%Zn-75 at% Cu alloy or a Sn-10 at% Zn-84 at% Cu alloy), the

Cu–Sn phase can be compounded in this Sn–Zn–Cu alloy without forming the Cu–Zn

phase [17].

Considering the stoichiometric ratio of CZTSe, the proportion of metallic components

is Cu:Zn:Sn = 2:1:1, thus the alloy can be written as a Sn-25 at% Zn-50 at% Cu alloy: Sn

and Zn concentrations are 25 at%, and Cu concentration is 50 at% in this alloy. This

composition indicates the coexistence of Cu6Sn5 and CuZn together with liquid Sn in

accordance with this phase diagram [17]. On the basis of the composition of the Sn-25 at%

Zn-50 at% Cu alloy, if the Cu concentration decreases and the proportion of Sn to Zn

remains the same (eg, Cu:Zn:Sn = 1:1:1), then Cu6Sn5 decomposes and forms liquid Sn

together with a Cu–Zn phase (such as CuZn or Cu5Zn8, etc). Contrary to this, if the Cu

concentration increases from the Sn-25 at% Zn-50 at% Cu alloy with the same proportion

of Sn to Zn (eg, Cu:Zn:Sn = 3:1:1), the increased Cu adheres not to CuZn but to Cu6Sn5

and forms a Cu-Sn phase (such as Cu6Sn5 or Cu3Sn). Consequently, when the Cu

concentration is 50–70 at% in the Sn-Zn-Cu alloy (eg, a Sn-20 at% Zn-60 at% Cu alloy),

the Cu-Sn alloys are observed together with CuZn. When the Cu concentration is higher

than 70 at% in the Sn-Zn-Cu alloy, such as a Sn-10 at% Zn-80 at% Cu alloy, CuZn

decomposes, and only Cu and the Cu-Sn phase are observed in this Sn-Zn-Cu alloy as

mentioned above.

2.1.2 The selenium–metal phase diagram

2.1.2.1 The Cu–Se thermodynamic system

At room temperature, the Cu-Se alloy can form six kinds of phases together with pure

Cu and Se, namely, α-Cu2Se, α-Cu2–xSe, β-Cu2–xSe, Cu3Se2, CuSe and CuSe2, depending

on the Se (or Cu) concentration [18]. As the temperature increases, the peritectic

decomposition occurs in Cu3Se2, CuSe2 and CuSe at different temperatures: Cu3Se2

decomposes into Cu2–xSe and CuSe at 113 °C, CuSe2 decomposes into Cu2–xSe and liquid

Se at 332 °C, and CuSe decomposes into Cu2–xSe and liquid Se at 379.7 °C [18]. The

temperatures of these phase decomposition are termed ‘peritectic decomposition

temperatures’.

When the Cu2–xSe or Cu2Se phases are compounded at room temperature by high

concentration of Cu in the Cu–Se alloy, these alloys do not decompose into other phases

2.2 Ellingham diagram for Se

5

unless the proportion of Se in the alloy changes.

2.1.2.2 The Zn–Se thermodynamic system

On a basis of the phase diagram for Zn–Se [19], ZnSe can be formed at room

temperature and coexists with pure Zn or Se, depending on the concentration of Se (or Zn).

As the Se concentration approaches 50 at%, the amount of ZnSe increases together with

the decrease in the amount of Se or Zn. When the temperature rises, no transformation

occurs, and no other phases in the Zn–Se alloy forms. Accordingly, only the transition of

pure Se or Zn from a solid to a liquid phase at those melting points occurs upon heating of

the sample.

2.1.2.3 The Sn–Se thermodynamic system

In the Sn–Se system, two kinds of intermediate compounds, SnSe and SnSe2, can be

formed at room temperature, depending on the Se (or Sn) concentration [20]. These two

Sn–Se alloys do not transform into other phases up to 628 C while pure Sn and pure Se

in the Sn-rich region (0–50 at% Se) and Se-rich region (~66.7–100 at% of Se) respectively

melt at ~231 and ~221 C [21]. Accordingly, SnSe and SnSe2 coexist in the Sn–Se alloy

until 628 C when the Se concentration is between 50 at% and 66.67 at%.

2.2 Ellingham diagram for Se

An Ellingham diagram for Se [22] is a graph that indicates the stability of binary

selenide depending on the temperature. The stability of each phase in this diagram is

determined by the Gibbs free energy (∆G), which is also termed ‘free energy’. This ∆G is

a numerical value for the preference for a reaction between one metallic element and the

Se in this case. A lower ∆G value indicates easier formation of its phase. In general, the

Ellingham diagram presents the metal which reacts with oxygon more easily. However,

this study needs the Ellingham diagram for selenium and not for oxygeon. Thomas B.

Reed has constructed the Ellingham diagram for selenium [22]. J.J. Scragg et al. also

calculated the free energy of formation of compounds in the metal–selenide binary

systems [23], and the calculated free energy shows the same tendency with the Ellingham

diagram built by Reed [22]. According to the Ellingham diagram, Zn is the metal which

has greatest tendency to form selenide, while Cu is the metal with the least tendency.

Therefore, here the Ellingham diagram for selenium [22] will be used to compare the

formation sequence of binary selenides.


6

2.3 Crystal structure data for the observed phases

As shown above, intermetallic phases in the alloy convert into other phases with

different crystalline structure or different unit cell dimensions as the temperature rises.

Table 2.1 displays information on the crystal structures of all phases used in this study.

Data here are obtained from the International Centre for Diffraction Data (ICDD) [24].

Table 2.1: The information on the references of all phases which are used for the diffraction analysis in this study.

Phase Space group Lattice parameters [Å] ICDD # (Temp.)

Mo Im-3m a = 3.147 97-064-3957 [25]

Cu Fm-3m a = 3.615 97-004-3493 [26]

Zn P63/mmc a = 2.665, c = 4.947

γ = 120° 97-005-2543 [27]

Sn I41/amd a = 5.831, c = 3.181 97-010-6072 [28]

Se P3121 a = 4.368, c = 4.958

γ = 120° 97-004-0018 [29]

Cu6Sn5 C2/c a = 11.022, b = 7.282, c = 9.827

β = 98.84° 97-010-6530 [30]

Cu3Sn Pmmn a = 4.772, b = 5.514, c = 4.335 97-015-0503 [31]

Cu41Sn11 F-43m a = 17.98 00-030-0510 [32]

Cu0.7Zn2 P-6 a = 4.275, c = 2.590

γ = 120° 97-010-3153 [33]

Cu5Zn8 I-43m a = 8.878 97-000-2092 [34]

CuZn Pm-3m a = 2.959 97-005-6276 [35]

Cu0.7Zn0.3 Fm-3m a = 3.584 03-065-9062

Cu2Zn (Cubic) a = 7.735 00-058-0457 [36]

CuSe

Cmcm a = 3.952, b = 6.962, c = 17.235 97-008-2330 [37]

P63/mmc a = 3.98, c = 17.254

γ = 120° 97-008-2331 [38]

Cu3Se2 P-421m a = 6.406, c = 4.279 97-001-6949 [39]

CuSe2 Pnnm a = 5.103, b = 6.292, c = 3.812 97-002-5717 [40]

a = 5.024, b = 6.194, c = 3.745 00-019-0400

Cu2–xSe F-43m a = 5.739 00-006-0680

Cu2Se Fm-3m

a = 5.787 97-004-1141 [41]

a = 5.871 97-010-3096 [42]

F23 a = 5.816 97-005-9955 [43]

ZnSe F-43m a = 5.633 97-004-1983 [44]

SnSe

Pnma

a = 11.501, b = 4.153, c = 4.445 97-005-2425 [45]

a = 11.559, b = 4.181, c = 4.429 97-005-0545 [46]

a = 11.571, b = 4.190, c = 4.419 97-005-0546 [46]

2.4 X-ray powder diffraction analysis

7

SnSe a = 11.589, b = 4.201, c = 4.409 97-005-0547 [46]

SnSe2 P-3m1 a = 3.795, c = 6.132

γ = 120° 97-065-1910 [47]

Cu2SnSe3 F-43m a = 5.684 03-065-4145

Cu2ZnSnSe4 I-42m a = 5.688, c = 11.338 97-009-5117 [48]

2.4 X-ray powder diffraction analysis

The phase compositions of samples are investigated by X-ray powder diffraction

(XRD). XRD is an analytical technique which reveals compounded phases in the sample

by observing diffracted X-ray beams from the sample. Because each phase has its own

unique unit cell dimensions (as discussed in section 2.3), the detected position and

intensity of the diffraction vary with the alloy phases. The principle of diffraction analysis

is explained in this section on a basis of references [49-51].

There are three methods for describing XRD phenomena from different points of view:

Laue equations, Bragg’s law, and Ewald sphere construction. Max von Laue approached

the analysis of diffracted beams in each axis of three-dimensional coordinates. The Bragg

analogy reduces these three dimensions into two dimensions for the explanation of this

phenomenon so that the experimental diffraction effects can be visualized as described in

Figure 2.1. When the monochromatic X-ray beam hits the crystal, the crystallographic

planes reflect this beam in accordance with the following equation (1) which is termed

“Bragg equation”:

nλ = 2dhklsinθ (1)

Here, n is integer, λ is a wavelength of X-ray beam, and dhkl is a distance between two

(hkl) planes of crystal: (hkl) denotes Miller indices for the orientation of a crystal plane.

However, because the crystal planes in real space are inherently three dimensional, the

Figure 2.1: The diffraction effect of X-ray beam on the (hkl) planes. The diffracted beam can be observed only when the distance ABC is equal to an integer multiple of a wave length λ of primary X-ray beam. This condition leads the Bragg’s law.


8

analysis of the real diffracted beam by this Bragg equation is frequently formidable.

The Ewald construction is the most useful method for describing and explaining

diffraction phenomena, as it introduces the concept of the ‘reciprocal lattice’ [49]. This

interpretation allows not only an understanding of the three-dimensional approach but also

the schematic approach, which incorporates Bragg’s law. The approach to diffraction

analysis focusing on Ewald sphere construction is explained in this section.

Reciprocal lattice

The reciprocal lattice [49, 50], or reciprocal space, is a new concept for the

observation of crystal structures from a different perspective. Its main concept is the

expression of an oriented crystalline (hkl) plane as a vector d*hkl. The direction of this

vector d*hkl is perpendicular to the (hkl) plane, and the magnitude of it is literally a

reciprocal number of the distance dhkl between two parallel (hkl) planes in the crystal.

Therefore the reciprocal d* vector can be defined as in equation (2).

|d*hkl| = 1/dhkl (2)

d*hkl = ha* + kb* + lc*

Here, a*, b* and c* are the basis vectors for the reciprocal lattice. They correlate with the

basic vectors of the crystal lattice, a, b and c as described in following equations (3):

𝒂*=𝒃×𝒄

𝒂∙(𝒃×𝒄), 𝒃*=

𝒄×𝒂

𝒃∙(𝒄×𝒂), 𝒄*=

𝒂×𝒃

𝒄∙(𝒂×𝒃) (3)

The diffraction detected from the single crystal represents the reciprocal lattice. That

is, the concept of this reciprocal space is not an imaginary space but a different coordinate

system for a crystal in real space. Furthermore, the observable lattice point, which is the

detectable diffraction along with the Miller indices of a crystal (hkl) plane, may be

predicted by Ewald sphere construction.

Ewald sphere construction

The Ewald sphere [50, 51] is the condition for diffraction of an X-ray beam. It has a

radius of 1/λ, and its centre is the crystalline sample, as described in Figure 2.2. When the

X-ray beam (S0) comes from the left side of the crystal sample into this Ewald sphere, the

point O which is diametrically opposite on the surface of the Ewald sphere from this

incoming X-ray beam becomes an origin of the reciprocal lattice: here, the length of S0 is

a radius 1/λ of this Ewald sphere. Based on this origin O of the reciprocal lattice, when

another reciprocal lattice point hkl (d*hkl) is also lying on the surface of the Ewald sphere,

2.4 X-ray powder diffraction (XRD) analysis

9

the incident X-ray beam is reflected to the [hkl] direction. In the case of Figure 2.2, the

diffraction S1 emerges from the crystal sample because the reciprocal lattice point 110 is

on the Ewald sphere together with the origin O. Accordingly, the length between the 000

and 110 points in the reciprocal space (|d*110|) denotes the reciprocal number of the

distance |d110|, in accordance with equation (2).

Figure 2.2: The correlation between Ewald sphere and diffraction together with the reciprocal lattice. When the reciprocal lattice point 110 (d*110) is lying on the surface of Ewald sphere together with the origin O of a reciprocal lattice, the (hkl) plane of crystal can diffract the incident X-ray beam (S0) to the direction of S1.

When the vector S0 moves to the right side of crystal, i.e. the vector S0 from the crystal

to the origin O of the reciprocal lattice as described in the left side of Figure 2.3, S0 forms

a triangle COP together with S1 and d*hkl, making an angle of 2θ with S1. This triangle

Figure 2.3: The relation between Ewald sphere construction and Bragg’s law. S0, S1, and d*110 form a triangle COP (left), and the Bragg equation can be induced by a right-angled triangle COH which is a half of triangle COP (right).


10

COP can also induce the Bragg equation by division of it into two right-angled triangles

as described in the right side of Figure 2.3. As mentioned above, the length of S0 and d*110

are 1/λ and 1/d110, respectively. Consequently, the length of a distance OH becomes

1/2d110 because the distance OH is a half of distance OP, and the value for sinθ denotes

the equation (1).

The above diffraction phenomena for a single crystal have been explained. However,

the samples for this study are a polycrystalline. For this reason, the X-ray beam is

diffracted to all possible directions while maintaining an angle of 2θ with respect to the

axis of the incoming beam. Consequently, a right circular cone is formed as described in

Figure 2.4, as if the vector S1 rotates about an axis of the vector S0. Thus, the apex of this

cone is the sample, and the aperture of this cone is 4θ. The circular ring on the base of this

cone indicates the diffracted beam from the polycrystalline sample and is termed ‘Debye–

Scherrer ring’ or ‘Debye ring’ [50, 51]. This Debye ring is detected and used for phase

analysis in this study.

Figure 2.4: The polycrystalline sample diffracts the X-ray beam as a circular ring which is termed ‘Debye-Scherrer ring’ or ‘Debye ring’, forming a shape of righ-circular cone. The aperture of this cone is 4θ.

3.1 Preparation of the Cu2ZnSnSe4 precursor

11

3. Experimental Details

3.1 Preparation of the Cu2ZnSnSe4 precursor

To observe the effect of the stacking sequence of elemental layers, various sequences

with different numbers of metallic components on the Mo-coated polyimide foil are

prepared. The polyimide foil (75 μm thickness) is used as a substrate for penetration of X-

ray beams through all samples in time-resolved in situ XRD measurements. The Mo layer

for all samples on the polyimide foil is also prepared to allow better adjustment of 2θ

scale of diffractions. Accurate angles of other diffractions may be derived from the Mo

reflection at 40.5°. The Mo-coated polyimide foil is prepared by Saint-Gobain Research

(SGR; France) by using a sputtering process. Afterwards, the components of CZTSe for

the metallic layers are deposited by sputtering in Avancis and by evaporation of a Se layer

in SGR and i-MEET. Table 3.1 describes the stacking sequence and the composition of

each sample along with its sample number. Samples are classified into three sections

according to number of metallic elements: one-, two- and three-metal samples. In all

Table 3.1: The composition of samples together with the sequence of elemental layers and the sample numbers. All samples can be divided into three sections depending on the number of deposited metals in its precursor, and this metal number is put on the sample number. The elemental layers are prepared with different sequences of elemental layers although the same metals are prepared in the sample.

Elemental

components

Sequence of

elemental layers

Elemental composition Sample

number Cu : Zn : Sn : Se

One-metal

sample

Zn–Se Mo/Zn/Se : 1 : : 1.3 Sample #1-1

Sn–Se Mo/Sn/Se : : 1 : 1.2 Sample #1-2

Two-metal

sample

Cu–Sn-Se Mo/Cu/Sn/Se 2 : : 1 : 3.3 Sample #2-1

Mo/Sn/Cu/Se 2 : : 1 : 3.3 Sample #2-2

Cu–Zn–Se

Mo/Zn/Cu/Se 2 : 1 : : 2.2 Sample #2-3

Mo/Cu/Zn/Se 2 : 1 : : 2.6 Sample #2-4

Mo/Cu/Zn/Se 1.3 : 1 : : 3.0 Sample #2-4a

Zn–Sn–Se Mo/Sn/Zn/Se 1 : 1 : 3.3 Sample #2-5

Mo/Zn/Sn/Se 1 : 1 : 3.3 Sample #2-6

Three-metal

sample Cu–Zn–Sn–Se

Mo/Cu/Sn/Zn/Se 1.8 : 1.2 : 1 : 5.3 Sample #3-1

Mo/Zn/Sn/Se/Cu 1.8 : 1.2 : 1 : 5.3 Sample #3-2

Mo/Zn/Sn/Cu/Se 1.8 : 1.2 : 1 : 5.3 Sample #3-3

Mo/Se/Cu/Sn/Zn 1.8 : 1.2 : 1 : 5.3 Sample #3-4

Mo/Zn/Cu/Sn/Se/Cu 1.8 : 1.2 : 1 : 5.3 Sample #3-5

Mo/Cu/Zn/Cu/Sn/Se 1.8 : 1.2 : 1 : 5.3 Sample #3-6


12

samples, Se is deposited at an excess stoichiometric ratio than CZTSe because it can

evaporate before combining with other elements during annealing.

The metal composition of the one- and two-metal samples is adjusted at the

stoichiometric proportion of CZTSe. Only the sample #2-4a is prepared with Cu-poor

composition for better understanding of reaction process between Cu and Zn. The three-

metal samples are differently prepared from the stoichiometric proportion of CZTSe as

described in Table 3.1, because the high efficiency CZTS(e)-based solar cell generally

consists of Cu-poor and Zn-rich composition [3, 52]. The metallic elemental components

of all samples are confirmed by X-ray fluorescence (XRF) at Avancis.

3.2 Characterisation techniques

Two kinds of XRD measurements are used to observe the alloy phases of samples:

time-resolved in situ XRD was used to observe the reaction path, and grazing-incidence

XRD (GIXD) was used to detect weak diffractions which may be hidden by the high

background signal of polyimide foil in the in situ XRD diffractogram. Because of the

similar diffraction angles of different alloy phases, Raman spectroscopy is also used to

confirm this compound.

3.2.1 Time-resolved in situ X-ray diffraction

Equipment for the in situ XRD measurement may be classified into three main

divisions: X-ray generator, sample chamber and detector parts. This technique is

performed as follows. A monochromatic X-ray beam from a rotating anode generator

passes through the sample which is mounted in the vacuum sample chamber. It then

reaches a charge coupled device (CCD) area detector as a ring diffraction pattern which is

termed ‘Debye-Scherrer rings’ or ‘Debye rings’. This measurement is performed during

heating of the sample, in accordance with the annealing process, as described below. After

the measurement, all detected Debye rings are converted into two-dimensional graphs of

intensity versus 2θ diffraction angles by using Fit2D program for the phase analysis. The

conditions for each part are described separately below. Information on the in situ XRD

set-up is described more detail in the dissertations of Hergert and Jost [53, 54].

X-ray generator

The X-ray beam is generated by a rotating copper anode operated at 44 kV and 75 mA

and is adjusted by two mirror optics to produce a strong Cu Kα radiation (λ = 1.5418 Å).

In front of these two-mirror optics, a vacuum cylinder is equipped to reduce air absorption


13

before hitting the sample. A 0.5 mm diameter slit is installed in front of this cylinder to

obtain a distinct diffraction. After passage through these apparatuses, the monochromatic

X-ray beam hits the sample mounted in the vacuum chamber perpendicularly.

Sample chamber

The sample chamber is prepared to prevent sample oxidation during annealing by

creating a vacuum state. Inside the sample chamber, a sample holder is fastened together

with a heater and a thermometer. As described in Figure 3.1, the pBN heater and the

Pt100 resistance thermometer are respectively fixed onto and beneath the sample holder

by two stainless-steel plates and two screws. To allow penetration of the X-ray beam, all

materials (sample chamber, heater, sample holder and two stainless plates) have a

concentric hole at the centre. Because the X-ray beam disperses after hitting the sample,

some parts have different diameter for the hole at the centre: the heater, sample holder,

and front-plate have 4.75 mm diameter, a back-plate with 25 mm diameter, and a front

and back sample chamber with 50 mm diameter. In the case of sample chamber, the holes

are covered by polyimide foil to make a vacuum inside the chamber.

Because of the low pressure during measurement (~0.1 Pa), the Se layer on the sample

Figure 3.1: Schematic diagrams of the sample chamber [55]. The sample holder (see Figure 3.2) is mounted by stainless steel and two screws in the middle of the sample chamber. The polyimide foils are transparent so that the X-ray beam can pass through it, and two beam stops behind the sample chamber and in front of the CCD area detector leads the well-defined diffractions.


14

Figure 3.2: Schematic diagrams of the sample holder [55]. The sample holder is made up of Cu and consists of concave and convex screws. Two pieces of sample are sandwiched between two polyimide foils as a face-to-face position and pressed by these screws.

easily evaporates from the sample. To prevent this evaporation, two pieces are taken from

one sample and are clamped by a sample holder, as described in Figure 3.2. The elemental

layer sides of the two pieces are in contact with each other, and the substrate sides are

separate. By using two pieces for one sample measurement, the diffraction intensity

becomes sufficiently strong to allow analysis of the alloy phases. For the same reason, the

sample pressure increases during the in situ XRD measurement because of vapourisation

of Se and/or SnSe [5, 56, 57]. Therefore, the sample sometimes ruptures under high

pressure before the temperature reaches 550 C. The use of two beam-stops behind the

sample chamber and in front of the detector (see Figure 3.1) also facilitates the analysis by

reducing the background signal effectively. These well-defined diffractions are collected

by a CCD detector, while the sample is heated in accordance with the annealing process.

Annealing process

The sample is heated from 30 to 550 C at a rate of 0.5 C/s, and then the temperature

is maintained at 550 C for 5 min. After the heating program, the sample cools down

naturally. All processes are performed at a pressure of ~0.1 Pa, and the diffractions are

steadily detected by a CCD area detector until the sample temperature is again reduced to

below ~300 C. This annealing process is performed by using a proportional-integral-

derivative (PID) controller (Eurotherm 2704) which is separately installed and connected

to the pBN heater in the sample chamber (see Figure 3.1). After the measurement, this

measured temperature must be recalculated on the basis of the Se melting points and/or

the peritectic decomposition temperatures of alloys which are already revealed in the alloy

phase diagrams. This is because the pBN heater does not apply heat directly to the sample

to allow penetration of X-ray beams through sample.


15

Detector

During annealing of the sample, the CCD area detector collects the Debye–Scherrer

rings dispersed from the sample behind the sample chamber, as described in Figure 3.1.

The surface of this detector is also mounted perpendicular to the incoming X-ray beam;

hence, the ring diffraction patterns can be detected, as shown on the left side of Figure 3.3.

However, it is necessary to convert these ring patterns into a one-dimensional graph of

intensity versus 2θ angle to analyse these diffractions by comparison with the references

of alloys (see section 2.3). The one-dimensional graph on the right side of Figure 3.3 is

the converted diffractogram from the ring diffraction patterns on the left side of Figure 3.3.

This conversion is performed by Fit2D program based on the position of the incoming

beam at 0° and the Mo reflection ring at 40.5°. The high background signal at around 12–

30° on the right side of Figure 3.3 is caused by the use of polyimide foil as a substrate.

Figure 3.3: Diffracted Debye–Scherrer rings obtained by the CCD area camera (left) and the converted graph of intensity versus 2θ degree (right). The conversion from ring diffraction patterns (left) into the one-dimensional graph (right) is performed by Fit2D program on a basis of the position of the incident X-ray beam at 0° and the Mo diffraction ring at 40.5°.

One diffractogram is recorded for 22.5 s (20 s illumination time and 2.5 s read-out

time), and a total 80 diffractograms are taken during one in situ XRD measurement until

the sample temperature decreases again from 550 C to below ~300 C. While the sample

temperature decreases from 550 °C to below ~300 °C, the obtained diffractions suggest no

phase transition in the alloy. Thus, only 60 diffractograms which are taken during

increasing the sample from 30 to 550 °C are enough to present the alloy reactions of the

sample.

Figure 3.4 on the left side, three-dimensional (3D) graph, displays 60 diffractograms

in a measuring sequence. Each diffractogram represents the compounded alloy phases at a

certain range of temperature. Because recording of one diffractogram takes 22.5 s as


16

mentioned above and because the temperature rises at a rate of 0.5 C/s, one diffractogram

can indicate the temperature of ~11 C. Consequently, this 3D graph presents the

transition of the diffractograms with respect to temperature. The right side of Figure 3.4,

two-dimensional graph (2D), is identical to the left side of Figure 3.4, but the intensity is

marked by colour instead of being presented on the z axis. One advantage of this colour-

coded 2D graph is the observable hidden diffractions behind the diffractions at high

intensity: the diffraction at ~29°, just behind the Se reflections (marked by red arrow),

cannot be observed on the 3D graph, whereas it can be observed on the colour-coded 2D

graph.

Figure 3.4: The collected 60 diffractograms while sample is annealed from room temperature to 550 °C for one in situ measurement. These diffractograms are displayed on a 3D graph (left) and a colour-coded 2D graph (right). Each diffractogram represents the compounded alloy phases at a certain range of temperature because recording of one diffractogram takes 22.5 s during a temperature rise of ~11 °C. One reflection which is marked by red arrow behind the Se reflections cannot be clearly observed on the 3D graph, whereas it can be observed on the 2D graph.

3.2.2 Grazing Incidence X-ray Diffraction

Some of precursors are measured by GIXD as an ex situ XRD measurement to verify

trace amounts of alloys in the precursor. Because occasionally one diffraction peak can

denote several alloy phases, these phases can be distinguishable by other small diffraction

peaks. However, when a trace amount of a related compound is included in the sample,

these small diffraction peaks are hidden in a high background signal which is produced by

polyimide foil. For this reason, the incident beam angle (ω) is fixed at 2° to reduce the

background signal by emitting the X-ray beam mostly on the surface of the film. It is deep

enough to allow detection of the Mo diffraction. A Philips X’pert Pro MPD powder


17

diffractometer operated at 40 kV and 35 mA is used for this measurement. Diffractions

are taken by a mini-prop detector mounted on the goniometer at a step size of 0.01° and an

exposure time of 20 s per step.

3.2.3 Raman spectroscopy

Raman spectroscopy is performed by Avancis to determine the presence of ZnSe in

the CZTSe film after the in situ measurement of the three-metal samples. Because the

main three diffractions of CZTSe have Bragg angles similar to those of the diffractions of

ZnSe and CTSe, it is necessary to investigate the samples by Raman spectroscopy.

Furthermore, the blue excitation laser is more effective at detecting ZnSe than the green

excitation laser [58]. On the basis of this result [58], the 488 nm blue laser with a spot size

of 1 μm is used. Area integration and point measurements are performed by a Horiba

Jobin Yvon LabRaman spectrometer at a spectral resolution below 0.5 cm−1

.

4. Results and Discussion

18

4. Results and Discussions

4.1 Investigation of single and binary metal systems with Se

Before starting the investigation into the formation of kesterite Cu2ZnSnSe4 (CZTSe)

film by using four components, the tendency of the reaction between Cu, Zn, Sn and Se is

determined in this section by observation of the reaction between Se and one or two

metallic elements. This tendency will be used to examine the reaction between four

components for a CZTSe formation in section 4.2.

4.1.1 Reaction path of a single metal with Se

In this section, selenisation of each metallic component is observed by using one-

metal samples. The reaction between Cu and Se, which is already revealed by Jost [54],

can be confirmed by the results for samples #2-2 (Mo/Sn/Cu/Se) and #2-3 (Mo/Zn/Cu/Se)

in section 4.1.2. Therefore, the Cu–Se sample is excluded, and only the reaction paths for

Zn–Se and Sn–Se are studied in this section.

4.1.1.1 Reactions of Mo/Zn/Se

To observe the formation temperature of the ZnSe phase without interferences, only

Zn and Se layers are prepared on the Mo-coated substrate for sample #1-1 (Mo/Zn/Se). As

shown in Figure 4.1 at ~30 C, only the Zn structure (peaks a) is detected in the precursor.

The strong reflection at 40.5° denotes a Mo structure which is a basis for 2θ angle

correction for other reflections. Except for Zn and Mo reflections, no Se or ZnSe

reflections are detected at this temperature, in contrast to those in the Zn–Se phase

diagram. According to the Zn–Se phase diagram (section 2.1.2.2), ZnSe and Se structures

coexist if the ratio of Se to Zn is more than 50 at% [19]; this is the case for this sample

([Se]/[Zn] = 1.3). This result suggests that the Se layer is in an amorphous state after

evaporation.

From ~110 to ~220 °C, the Se layer crystallises while the Zn structure grows along the

[002] direction. As the temperature increases, the Se reflections which could not be

detected in the precursor become clearly observable and then disappear at ~220 C, its

melting point (221 °C). Meanwhile, two of peaks a (Zn) at 36.29° and 43.23° in Figure 4.1

shift to smaller Bragg angles: the peak at 36.29° shifts further than the peak at 43.23°.


19

Figure 4.1: Time–temperature evolution of powder diffractograms of sample #1-1 (Mo/Zn/Se) with colour-coded intensities (a.u.) of the Bragg reflections. Phases denoted by peaks are as follows: a: Zn, b: ZnSe.

These two reflections at 36.29° and 43.23° denote (002) and (101) planes, respectively. It

means that Zn structure grows along the z axis.

At ~290 C, ZnSe is compounded at low formation rate. As the peaks b (ZnSe) appear

at this temperature, peaks a (Zn) gradually weaken. The change in their intensities shows

that conversion of Zn into ZnSe does not occur at once, but rather slowly. This slow

transformation signifies slow penetration of liquid and/or a gaseous Se into the

crystallised Zn layer. While the upper side of Zn layer alloys with Se, the lower side of Zn

layers still maintains its structure due to the slow rate of ZnSe formation. For this reason,

Zn is detectable together with the ZnSe reflections.

The ZnSe formation temperature at ~290 C is much higher than its possible

formation temperature at 221 C, in accordance with the Zn–Se phase diagram [19].

Although Se melts at ~220 C in this measurement, Se could not react with Zn until

~290 C. This result implies that a certain amount of thermal energy is necessary to form

ZnSe when Zn is prepared as a pure Zn layer with a well crystalline structure.

Because ZnSe forms slowly, the Zn reflections (peak a) disappear at ~390 C, ~100 C

higher than the ZnSe formation temperature. This temperature of disappearance of Zn

reflections is lower than its melting point, 419 C, but it is reasonable considering the use

of Zn in ZnSe formation.

4.1.1.2 Reactions of Mo/Sn/Se

SnSe formation in sample #1-2 shows a result different from those of the formation in

sample #1-1. SnSe forms as soon as Sn melts, whereas ZnSe forms before Zn melts. On


20

Figure 4.2: Time–temperature evolution of powder diffractograms of sample #1-2 (Mo/Sn/Se) with colour-coded intensities (a.u.) of Bragg reflections. Phases denoted by peaks are as follows: c: Sn, d: SnSe, e: SnSe2, E: SnSe (d) and SnSe2 (e) [59].

the contrary, the beginning reaction of this sample #1-2 seems to be similar with that for

the sample #1-1, on the basis of its different metallic element.

Except for the Mo reflection, only Sn reflections (peaks c) are detected in the

precursor, as shown in Figure 4.2 at ~30 C. Here, the Se reflections are also not observed

for the same reason as that for sample #1-1, that is, the amorphous Se layer. One peculiar

observation at this stage is the tetragonal Sn structure oriented towards the [101] direction

relative to the Sn reference (ICDD #97-010-6072), as shown in Figure 4.3. The Sn layer

appears to be influenced by a Mo layer, which is oriented along the [110] direction, during

the sputtering process.

Figure 4.3: The comparison between X-ray diffractogram of sample #1-2 (Mo/Sn/Se) and reference Sn data from ICDD #97-010-6072. This diffractogram is one of diffractograms in Figure 4.2, which is measured at ~30 °C [59].


21

While Se crystallises from ~110 C and vanishes at ~220 C, Sn reflections (peaks c)

slightly and gradually weaken from ~120 C and suddenly disappear at ~230 C. The in

situ analysis (Figure 4.2) shows these diminishing peaks c at 30.64° and 32.03°, but the

peaks c at 43.88° and 44.91° are not observable because Se reflections overlap on the right

side of these peaks. This weakening Sn reflection implies that a crystallising Se layer

slightly influences the Sn structure. The temperatures of disappearance of these Se and Sn

reflections agree with the melting points for Se and Sn at 221 and 232 C, respectively.

After Sn melts at ~230 C, the SnSe phase (peak d and E) immediately forms from

liquid Se and liquid Sn. When the peaks c (Sn) disappear at ~230 C, one of them (at

~30.6°) intensifies again along with peak d, which represents SnSe. The peak at ~30.6°,

denoted by a capital E, indicates not only SnSe but also SnSe2 phases due to the same

Bragg angle of each main reflection. The phase represented by peak E may be

distinguishable from the additional reflections for SnSe at 37.78° (peak d, ICDD #97-005-

2425) and for SnSe2 at 14.43° (peak e, ICDD #97-065-1910). Thus, the beginning of peak

E indicates SnSe by the emergence of peak d from ~230 C.

Afterwards, SnSe2 (peak e and E) also alloys from ~270 °C. As mentioned in the

above paragraph, SnSe2 is distinguishable from SnSe by the additional peak at 14.43°.

This peak e (SnSe2) slowly emerges from ~270 C as peak d (SnSe) vanishes during

transformation of SnSe into SnSe2. Acording to the Sn–Se phase diagram [13], SnSe can

co-exist with SnSe2 if the Se concentration is between 50 and 65 at%. In the case of this

sample, the ratio of Se to Sn is 1.2 ([Se]/[Sn] = 1.2), ie, SnSe and SnSe2 could co-exist in

this stage. However, the distinguishable peak d (SnSe) is not detectable from ~310 °C

until the sample ruptures. This could be caused by the domination of SnSe2 phase over the

film. Accordingly, the weak reflection of SnSe produced by the small amount of this

phase seems to be veiled by a relatively high background signal relative to the weak

reflection.

When the sample ruptures at ~520 C, leading to a decrease in sample pressure, the

SnSe2 phase dealloys again into SnSe and Se at ~540 C. As mentioned in section 3.2.1,

these clamped two pieces of sample can rupture at high temperature by vapourisation of

Se and/or SnSe during measurement because the substrate is polyimide foil. This sample

rupture may be recognised from the shifting reflections, especially those of the Mo peak.

Figure 4.2 also presents the rapid shift of all reflections at one time together with the Mo

reflection at ~520 °C, which signifies the decrease in pressure during the measurement of

sample. After the shift of all reflections, peak e disappears and peak d appears again,

indicating the transition of SnSe2 into SnSe. Although peak e disappears at this stage, the

presence of a small amount of SnSe2 in the sample is also possible, in accordance with the

Sn–Se phase diagram [13]. However, the dominant phase in this sample at this stage is

SnSe. Dealloying of SnSe2 to SnSe after sample rupture implies that the sample pressure


22

can affect the formation of the Sn–Se compound in the film: higher pressure induces the

SnSe2 formation, and lower pressure causes the SnSe formation. Other studies [56, 57]

have revealed that high pressure is necessary to prevent the evaporation of SnSe as a gas,

but this high pressure can also induce the SnSe2 formation on a basis of this result.

4.1.2 Reaction path of two metals with Se

The reaction of Se with two metallic elements is observed to understand better the

correlation between Cu–Zn–Sn–Se elements in the alloy. The results show different

formation temperatures for same alloy, along with different reaction paths, depending on

the sequence of stacked layers in the precursor. It proves that the layered sequence affects

the alloy formation temperature more than does the elemental ratio and temperature as

indicated in the alloy phase diagram. By means of this observation, the characteristics and

the tendency for reaction between four components is examined in section 4.1.3.10 and

becomes the basis for an understanding of CZTSe formation in section 4.2.

4.1.2.1 Reactions of Mo/Cu/Sn with Se

The precursor of sample #2-1 (Mo/Cu/Sn/Se) consists of Sn, Cu6Sn5 and metallic Cu

phases together with an amorphous Se layer. The two elements, as well as Sn (peak c) and

Cu6Sn5 (peaks f and Z), are clearly detected at ~30 C, as shown in Figure 4.4. However,

metallic Cu is unclear because of the overlap of the Cu reflection with that of Cu6Sn5

reflections at ~43.3 (peak Z). Considering the components of the precursor and its ratio

of elemental Cu to Sn ([Cu]/[Sn] = 2), metallic Cu obviously exists in the precursor;

therefore, peak Z also includes the Cu reflection. As both Cu and Cu6Sn5 phases, as well

as many other phases, have a strong reflection near peak Z, this peak is present in most

observations in this study, denoting various phases. The first and second phases for this

peak Z are Cu and Cu6Sn5, and the third phase is Cu3Sn which evidences outward

diffusion of Cu in this sample.

As the temperature rises, Cu3Sn and CuSe appear at ~180 and ~200 C, respectively,

because of outward diffusion of Cu. While the Se layer crystallises at ~110 C and melts

at ~220 C, Cu outwardly diffuses through the Sn layer to the Se layer. In Figure 4.4,

peaks c (Sn) and f (Cu6Sn5) gradually weaken as the Se reflections become detectable.

Meanwhile, peaks g (Cu3Sn) slowly emerge from ~180 C, and then peak h (CuSe)

appears at ~200 C as soon as peak c (Sn) vanishes. Here, the small peaks g match closely

the reference data for Cu3Sn (ICDD #00-015-0503), including the strongest reflection at

~43.2° (peak Z). Peak h also corresponds well with reference data for CuSe (ICDD #97-


23

Figure 4.4: Time–temperature evolution of powder diffractograms of sample #2-1 (Mo/Cu/Sn/Se) with colour-coded intensities (a.u.) of the Bragg reflections. Phases denoted by the peaks are as follows: c: Sn, d: SnSe, f: Cu6Sn5, g: Cu3Sn, h: CuSe, i: l-CuSe2, j: s-CuSe2, k: Cu2SnSe3. Reflections for CuSe2 change from peak i to peaks j because the unit cell size of this CuSe2 phase is reduced. Here, l-CuSe2 and s-CuSe2 indicate a relatively large and small unit cell for CuSe2, in accordance with references data from ICDD #97-002-5717 and ICDD #00-019-0400, respectively.

008-2330). Therefore, the transformation of Cu compounds can be described by the path

Cu6Sn5 → Cu3Sn → CuSe. This reaction path clearly confirms Cu diffusion from the

bottom to the top of the film, passing through the Sn layer according to the sequence of a

precursor, Mo/Cu/Sn/Se. Additionally, CuSe formation from the Cu3Sn phase confirms

that Cu has a stronger tendency to react with Se than with Sn: (i) Cu–Sn < Cu–Se.

At ~230 C, CuSe2 (peaks i and j) forms because of the sudden increase in the liquid-

phase concentration of Se due to liquefied Se at 221 C. This formation reaction accords

well with the Cu–Se phase diagram; as the Se concentration increases, more CuSe phase

alloys into CuSe2 when the Se concentration is between 50 and 66 at% [18]. In our result,

this CuSe2 phase has two types different unit cell sizes, according to the reference data

from ICDD #97-002-5717 for a larger unit cell (peak i, l-CuSe2), as well as another

reference data from ICDD #00-019-0400 for a smaller unit cell (peak j, s-CuSe2), as

described in Table 2.1 in section 2.3. As soon as the Se reflections disappear, a weak peak

i appears at 29.22° in a subsequent diffractogram. After an increase of ~10 C, peak i

shifts to larger Bragg angles at 29.76°, which is denoted by ‘j’. This change suggests that

CuSe2 with a large unit cell forms at the beginning because of the higher concentration of

Se and then becomes a smaller because of the steady penetration of Cu and/or the Se

diffusion into the film.

At ~250 °C, SnSe forms from Sn decomposed from Cu3Sn and Se diffusing from


24

CuSe2, resulting in complete decomposition of Cu3Sn and Cu6Sn5 at this temperature. This

may be deduced from the intensity variation of these peaks. As peaks g (Cu3Sn), f and Z

(Cu6Sn5) vanish at this temperature, peaks d (SnSe) emerge at ~30.5° and ~38° together

with the growing peaks h (CuSe). This means that all of the Cu from the Cu–Sn alloys

decomposes, forming a Cu–Se compound, and the remaining Sn becomes a liquid phase

because the temperature is higher than its melting point. Therefore, Sn can combine with

Se. At this point, the reflections of CuSe2 (peaks j) show the strongest intensity and then

gradually weaken from ~260 C, while other reflections of SnSe (peak d) and CuSe (peak

h) strengthen. This indicates that Se for the formations of CuSe and SnSe phases derives

from the CuSe2 phase; thus, the amount of CuSe also increases along with the amount of

SnSe:

CuSe2 + Sn(l) → CuSe + Se + Sn(l) → CuSe + SnSe (at ~260 °C) (1)

One interesting observation at this reaction stage is that the first binary selenide is not

SnSe but CuSe. Sn does not react with Se although the Sn layer is in contact with the Se

layer in the precursor and can undergo an immediate reaction with Se as shown in sample

#1-2 (Mo/Sn/Se). Cu instead of Sn reacts with Se when Cu permeates the film through the

Sn layer. This phenomenon indicates that Se has a stronger tendency to react with Cu than

with Sn: (ii) Sn–Se < Cu–Sn.

At ~290 °C, the formation of Cu2SnSe3 (CTSe) starts, and its formation rate increases

as all of CuSe2 decomposes into CuSe and Se at ~330 C. This change may be observed

from the different increase in rate of the intensity of peak k. As shown in Figure 4.4, peak

k (CTSe) emerges at ~290 C at ~27°, where is the left of peak h at ~28°, forming a

shoulder-like peak. At that time, peaks j (CuSe2) steadily weaken whereas peaks h have

unchanged intensities. This indicates the involvement of CuSe2 in CTSe formation at

~290 C, as described by equation (2). When the temperature reaches ~330 C, peak k

(CTSe) clearly emerges as evidenced by the increase in intensities. Here the temperature

at ~330 °C corresponds to the peritectic decomposition temperature of CuSe2 at 332 °C

[18]. Although the amount of CuSe increases with decomposition of CuSe2 into CuSe,

peaks h (CuSe) gradually weaken alongside peaks d (SnSe) at this temperature. It means

that the components for the CTSe formation changes from CuSe2 to CuSe at ~330 °C.

Addtionally, the rate of CTSe formation increases relative to that before reaction with

CuSe2 and SnSe, as described by equation (3). These two reactions with different rates of

CTSe formation suggest that the rate of CTSe formation depends on Cu–Se compounds.

The depending of the formation rate on the Cu–Se compounds is also observed with CIGS

in other studies [60].


25

2 CuSe2 + SnSe → Cu2SnSe3 + 2 Se (~290–330 °C): relatively slow (2)

2 CuSe + SnSe → Cu2SnSe3 (>330 °C): relatively fast (3)

Traces of Cu2Se (peak l) is shortly observed at around 380–400 °C. However, it seems

to be involved in the formation of CTSe, as evidenced by the disappearance of Cu2Se

reflection. While two diffractograms are taken during the increase of ~20 °C (from 380 °C

to 400 °C), the intensity of peak l weakens and this peak l soon vanishes as described in

Figure 4.4.

4.1.2.2 Reactions of Mo/Sn/Cu with Se

The precursor with an inverse order of Cu and Sn layers with respect to that of sample

#2-1 is prepared. Accordingly, the Cu layer for sample #2-2 is in contact with Se on the

upper part of the film: Mo/Sn/Cu/Se. Similar to the above sample, Cu6Sn5 (peaks f and Z)

and Cu (peak Z) are also detected in the precursor, as shown in Figure 4.5 at ~30 C. As

mentioned above, metallic Cu is apparently included in peak Z, in accordance with the

elemental ratio. Differences upon change of the sequence of elemental layers at this stage

are the weak reflection of CuSe (peak h) and the undetectable Sn reflections (peaks c in

Figure 4.5: Time–temperature evolution of powder diffractograms of sample #2-2 (Mo/Sn/Cu/Se) with colour-coded intensities (a.u.) of the Bragg reflections. Phases denoted by peaks are as follows: c: Sn, d: SnSe, f: Cu6Sn5, h: CuSe, i: l-CuSe2, j: s-CuSe2, k: Cu2SnSe3. Reflections for CuSe2 change from peak i to peaks j because the unit cell size of this CuSe2 phase is reduced. Here, l-CuSe2 and s-CuSe2 indicate a relatively large and small unit cell for CuSe2, in accordance with references data from ICDD #97-002-5717 and ICDD #00-019-0400, respectively.


26

Figure 4.6: GIXD diffractogram of sample #2-2 (Mo/Sn/Cu/Se) before heat treatment. This analysis is performed in the atmosphere at room temperature. Although Sn is deposited between Mo and Cu layers in the precursor, only a trace of Sn reflection is detected.

Figure 4.4). Because Cu is a reactive material, a small amount of Cu seems to combine

with Se during sample preparation. The presence of CuSe in the precursor of sample #2-2

proves the reactivity Cu in comparison with the absence of SnSe in the precursor in

sample #2-1. Another difference is the undetectable Sn reflections, which seem to be

caused by the Cu deposition on the Sn layer. In the sputtering process, the sputtering

element has enough energy to reach with other element, and, as mentioned above, Cu is

very reactive. Therefore, Cu combines with most of the Sn in the sample during sputtering

of Cu on the Sn layer. This is also verified by ex situ analysis described in Figure 4.6. To

confirm the presence of metallic Sn, the precursor is precisely analysed by GIXD. This

additional measurement detects traces of Sn reflections. It also indicates that most of the

Sn has already combined with Cu during sputtering, as expected.

While Se crystallises and melts at 110–210 C, CuSe grows with crystalline Se and

metallic Cu. As the temperature rises, a pure metallic Cu which has not combined with Sn

during sample preparation reacts with Se. In particular, this gradual formation of CuSe

from Cu is notably observed from the gradual weakening of peak Z together with the

strengthening of peak h (CuSe) when the Se reflections appear at ~110 C during Se

crystallisation. As mentioned above, peak Z denotes Cu and Cu6Sn5. In this case, however,

it is certain that the diminishing peak Z belongs only to metallic Cu because peak f

(Cu6Sn5) at ~30° does not weaken. The inverse change of intensities between peak Z and

peak h (CuSe) also confirms this observation. Because of the active phase transition, Cu +

Se → CuSe, the Se reflections diminish rapidly at ~200 C and disappear completely at

~210 C, which is lower than its melting point. This signifies that the reactivity of Cu is

high, as it combines with the crystallised Se layer, which does not occur in the case of Sn.

Sn reacts with liquid Se but not with crystalline Se (see section 4.1.2.1).

As Cu diffuses from Cu6Sn5 through the film, metallic Sn and l-CuSe2 emerge at ~180


27

and ~190 °C, respectively, along with the growth of CuSe. While CuSe (peak h) is formed

by a crystallized Se and a metallic Cu (peak Z), Cu dealloyed from Cu6Sn5 (peaks f and Z)

also reacts with Se, forming l-CuSe2 (peak i). Because Cu tends to diffuse outwardly, this

CuSe2 phase seems to form on the upper part of the film. Meanwhile, the remaining Sn

from Cu6Sn5 bonds together and crystallises as a tetragonal Sn structure (peaks c) on the

lower part of the film. Therefore Sn reflections which were not detected in the precursor

become observable at ~180 °C in Figure 4.5, and consequently Cu6Sn5 decomposes at

~190 °C. Here the decomposition temperature of Cu6Sn5 accords with the temperature for

the transformation from η’–Cu6Sn5 into η–Cu6Sn5 at 186–189 °C [13]. According to the

phase diagram for Cu–Sn alloy [13], these two phases have hexagonal structure although

both are different phases. However, the Cu6Sn5 reflections (peaks f and Z) in Figure 4.5

signify a monoclinic structure, according to ICDD #97-010-6530 data, which matches

closely peaks f and Z. That is, when monoclinic Cu6Sn5 changes into a hexagonal

structure at around 186–189 °C, decomposed Cu from the nomoclinic structure does not

adhere to Sn again to form a hexagonal η–Cu6Sn5 but instead adhere to Se, diffusing

outward. Because of the increase in Cu concentration on the Se layer, substantial amounts

of CuSe form together with l-CuSe2, which has a large unit cell. Figure 4.5 also shows the

rapid weakening and disappearance of peaks f and Z (Cu6Sn5) at ~180 and ~190 C,

respectively. Subsequently, peak i (l-CuSe2) and peaks c (Sn) appear, and peak h (CuSe)

reaches the highest intensity during measurement. The formation of l-CuSe2 instead of

SnSe at this stage also verifies that Se has a stronger tendency to react with Cu than Sn:

(iii) Sn–Se < Cu–Se.

At ~230 °C, the unit cell of CuSe2 shrinks simultaneously with the formation of SnSe

because of the Se diffusion through film. As soon as peaks d (SnSe) and j (s-CuSe2)

appear at ~230 °C, peaks c (Sn) and i (l-CuSe2) disappear and one of peaks h (CuSe) at

~27° weakens. This significant change in the diffraction patterns signifies that metallic Sn

may react with Se as the unit cell size of CuSe2 becomes smaller (l-CuSe2 → s-CuSe2).

This reaction process can be interpreted only by the Se diffusion from l-CuSe2 into the

film. For the same reason, the CuSe reflections (peak h) become weaker. The inverse

change in intensities between peaks j (s-CuSe2) and peaks d (SnSe) near this temperature

also demonstrates the Se diffusion form s-CuSe2 through film. As peaks j slightly weaken,

peaks d (especially at ~38°) gradually strengthen at 250–290 °C. This signifies the

increase in the amount of SnSe on the lower part of the film along with the decrease in the

amount of s-CuSe2 on the upper part of the film due to the Se diffusion into the film.

At ~290 C, CTSe forms from CuSe2 and SnSe at low rate of its formation. In Figure

4.5, the CTSe reflection (peak k) at ~27° is so faint that it is a bit difficult to distinguish

from the background signal. On the contrary, this reflection becomes clearly observable

and stronger as soon as CuSe2 (peaks j) transforms into CuSe (peaks h) at ~330 C: the


28

peritectic decomposition temperature of CuSe2 is 332 °C [18]. This different rate of CTSe

formation is similarly to the result for sample #2-1, as described in equation (2) and (3)

(see section 4.1.2.1).

At 380–400 °C, traces of Cu2Se are detectable from the weak reflections (peak l) as

the CuSe reflections (peaks h) disappear at ~380 °C. This temperature is well in

accordance with the peritectic decomposition temperature of CuSe at 379.3 °C [18]. This

compounding Cu2Se is soon integrated into the CTSe phase, similar to sample #2-1.

Therefore peak l is not observable above 400 °C.

The difference in results between sample #2-1 and #2-2 is undetectable Cu3Sn in

sample #2-2 because most of Sn already reacts with Cu during sputtering. For the same

reason, the reflections of Sn (peaks c) in sample #2-2 were not observed at ~30 °C but

appear at ~180 °C as soon as Cu diffuses outwardly from Cu6Sn5 to Se layer, resulting the

decomposition of Cu6Sn5 (peak f and Z) at ~190 °C.

4.1.2.3 Reactions of Mo/Zn/Cu with Se

A weak reflection for CuSe (peak h) and peak Z for sample #2-3 (Mo/Zn/Cu/Se) are

observed at an early stage of measurement (see Figure 4.7). As mentioned above, peak Z

may indicate not only elemental Cu and Cu6Sn5 but also elemental Zn, Cu5Zn8 and CuZn

because of the overlap of its strongest Bragg reflections with those of phases at 2θ values

of ~43.2°. As the Cu6Sn5 phase are recognisable from the additional peak at 30.15° in

previous results, the Zn and Cu5Zn8 phases are also distinguishable by additional weak

reflections at angles of 36.29° and 39.00° for the Zn phase (ICDD #97-005-2543) and at

angles of 34.98° and 37.89° for the Cu5Zn8 phase (ICDD #97-000-2092). Because these

additional peaks which make phases distinguishable have weak intensities, high

background signal produced by polyimide foil can obscure these peaks when the sample

includes those phases at trace amounts. For this reason, GIXD measurement is performed

to determine the presence of Zn and Cu5Zn8, which cannot be confirmed by in situ

analysis (Figure 4.7). The diffractogram taken from GIXD measurement confirms the

absence of Zn and Cu5Zn8, as shown in Figure 4.8. Therefore, it is clear that peak Z

indicates only the CuZn phase together with metallic Cu according to the elemental ratio

([Cu]/[Zn] = 2). It suggests that most of the Zn combines with Cu, forming CuZn during

sputtering Cu on the Zn layer. Thus, the precursor becomes having a sequence as follows:

Mo/CuZn/Cu/Se. This is similar to Sn in sample #2-2, which also shows the combination

of most of the Sn with Cu during sputtering (see section 4.1.2.2). The presence of CuSe in

the precursor also verifies the combination of Cu with Se during a deposition of Se layer

on the Cu layer, similar to the result for sample #2-2. According to observation of these

compounds in the precursor, Cu is the most reactive element relative to Zn or Sn because


29

Figure 4.7: Time–temperature evolution of powder diffractograms of sample #2-3 (Mo/Zn/Cu/Se) with colour-coded intensities (a.u.) of the Bragg reflections. Phases denoted by peaks are as follows: b: ZnSe, h: CuSe, i: l-CuSe2, j: s-CuSe2, l: Cu2Se, Z: Cu, CuZn and possibly Zn and Cu5Zn8.

Figure 4.8: GIXD diffractogram of sample #2-3 (Mo/Zn/Cu/Se) before heat treatment. This analysis is performed in the atmosphere at room temperature. Neither Zn nor Cu5Zn8 in the precursor is detectable by this ex situ analysis.

only a part of Cu reacts with Se during preparation, whereas Zn and Sn do not (see results

for samples #1-1 and #1-2).

Upon heating of the sample, an amorphous Se layer crystallises and then actively

reacts with Cu, forming CuSe at ~150 °C and l-CuSe2 at ~180 °C at high rate. These

reactions can be found in Figure 4.7 by change in intensities of peaks h, i and Z and Se

reflections. When peak h (CuSe) rapidly strengthens at ~150 °C, leading to the emergence

of additional CuSe reflections at ~26° and ~31°, Se reflections which appeared at ~110 °C

weaken. Simultaneously with this, the intensity of peak Z (Cu, CuZn) also suddenly


30

decreases. The change in intensities of these diffraction patterns intends the use of

metallic Cu and crystalline Se for the formation of CuSe at this temperature (~150 °C).

Afterwards, the weakened Se reflections completely disappear at ~180 °C as peaks i (l-

CuSe2) emerge. Here, the earlier disappearance of Se reflections relative to its melting

point is reasonable because of the rapid consumption of Se in the formation of Cu–Se

alloys. The formation of l-CuSe2 along with the vanishing Se reflections signifies the

steady increase in Cu concentration in the Se layer. Because pure metallic Cu elements,

which were not combined with Zn in the precursor (Mo/CuZn/Cu/Se), continuously

penetrates from bottom to top of the crystallised Se layer, l-CuSe2 forms at ~180 °C.

Meanwhile, the distribution of Cu through β’-CuZn is steadily changed at 150–220 °C,

forming Cu-rich and Zn-rich Cu–Zn alloys, because of the dezincification. While Cu–Se

phases are formed on the upper part of film, Cu steadily decomposes from a β’-CuZn

phase (peak Z) by dezincification and diffuses to the outward film where the Se layer is.

In fact, the peak Z at ~150 °C denotes only β’-CuZn because of the use of all metallic Cu

for the CuSe formation below ~150 °C. Afterwards, this peak Z gradually shifts to low

Bragg angles that may be interpreted in two ways: the growth of the β’-CuZn structure

due to the rise in temperature and the transformation of CuZn to the Cu-rich Cu–Zn alloy

which diffracts the X-ray beam at a Bragg angle lower than that for β’-CuZn, such as

Cu2Zn (ICDD # 00-058-0457). If the first interpretation (a growing structure due to a rise

in temperature) causes the shift of peak Z, the reflection of CuZn phase moves gradually

to low Bragg angles until the reflection disappears. However, the shift of peak Z stops at

~220 C and settles at the Bragg angle until it disappears; thus, the second interpretation,

the formation of Cu-rich Cu–Zn alloy, seems reasonable. According to a study on the

transformation from bcc to fcc phase in Cu–Zn alloy [61], the Zn component evaporates

from Cu–Zn alloy when sufficiently high temperature is supplied. Additionally, a Cu–Zn

phase diagram [8] implies decomposition of Cu from CuZn (see section 2.1.1.1). When

the Zn concentration is near 50 at%, a certain amount of CuZn converts into Cu5Zn8 and

Cu during a temperature increase from 227 to 468 C, as described by following equation:

8 CuZn → Cu5Zn8 + 3 Cu. These clearly proves the occurrence of dezincification in the

β’-CuZn phase. For this reason, Zn steadily and selectively leaches from the β’-CuZn

phase, and at the same time, the remaining Cu diffuses into the upper side of the CuZn

alloy, producing a Cu-rich Cu–Zn phase. Therefore, peak Z steadily shifts to low Bragg

angles becasue of the gradual compounding of Cu-rich CuZn alloys in this stage. In other

words, the shift of peak Z represent a decline in Zn concentration from the bottom to the

top of the Cu–Zn structure, similar to a result of other study [62].

The shifting of peak Z stops the Bragg angle at ~220 C. The halted shift of peak Z

signifies that the formation of Cu-rich Cu–Zn alloy on the upper side of Cu–Zn alloy has

stopped. This indicates that the Cu dealloyed from the Cu–Zn alloy cannot diffuse into


31

Cu–Se alloy but remains in the Cu–Zn alloy, maintaining the Cu concentration on the

upper side of Cu–Zn alloy. The restriction of Cu diffusion into the Se element in this

sample is the converse of a result for sample #2-1 (Mo/Cu/Sn/Se). It signifies that, once

Cu–Zn alloy forms in the precursor, the combination between Cu and Zn is stronger than

between Cu and Se: (iv) Cu–Se < Cu–Zn.

At the same temperature (~220 °C), Se diffuses from Cu–Se alloy into the film. This

reaction is evidenced by decreasing intensities of peaks h (CuSe) and peak i (l-CuSe2)

together with the emergence of peaks j (s-CuSe2). These inverse changes in intensities

also imply the transformation of l-CuSe2 into s-CuSe2 by means Se diffusion into the film,

similar to the result for sample #2-2 (Mo/Sn/Cu/Se).

At ~290 °C, ZnSe (peak b) starts to form from Se diffusing from CuSe2 (peak j) and

Zn decomposed from the Cu–Zn alloy (peak Z). Simultaneously with this reaction, the

amount of CuSe increases by Se diffusion from CuSe2 (peak j). The consumption of

CuSe2 instead of CuSe in the formation of ZnSe may be verified by the weakening of

peaks j (s-CuSe2) and the emergence of the faint peak b (ZnSe). At the same time, one of

additional CuSe reflections near ~31° (peak h) emerges. They signify that ZnSe forms

together with CuSe through equation (4). Additionally, peak Z, which denotes different

proportions of Cu–Zn alloys, vanishes because of the use of Zn from the Cu–Zn alloy for

the ZnSe formation. The formation temperature of ZnSe phase also matches closely the

result for sample #1-1 (Mo/Zn/Se). Apparently, the Cu–Se alloy does not interrupt the

reaction between Se and Zn but instead facilitate, similar to the reaction of Se and Zn in

sample #1-1 (Mo/Zn/Se).

CuSe2 + (dealloying) Zn → CuSe + ZnSe (290 C <) (4)

As CuSe2 dealloys into CuSe at ~330 C, the amount of ZnSe increases a bit.

Accordingly, as soon as peaks j disappear, the peak b at ~27° strengthens, as shown in

Figure 4.7.

At ~380 °C, which is near the peritectic decomposition temperature of CuSe, all of the

CuSe (peaks h) dealloys into Cu2Se (peaks l), and the amount of ZnSe (peak b) suddenly

increases because of residual Se from the decomposition of CuSe to Cu2Se: 2 CuSe →

Cu2Se + Se. This result includes that the amount of ZnSe may increase when the Cu–Se

alloy decomposes.

4.1.2.4 Reactions of Mo/Cu/Zn with Se

The reaction path for sample #2-4 (Mo/Cu/Zn/Se) described in Figure 4.9 seems to be

simple in comparison with that for sample #2-3 (Mo/Zn/Cu/Se). Only two kinds of


32

reflections, peak Z and Se, are observed until 360 C in the absence of any reflections of

Cu–Se compounds. Considering the result for sample #2-1 (Mo/Cu/Sn/Se), which clearly

shows Cu diffusion through the Sn layer with the formation of Cu–Sn and Cu–Se

compounds, the Zn layer appears to block the outward diffusion of Cu, as evidenced by

the undetectable reflections of Cu–Se alloys. To investigate in detail the alloys

corresponding to peak Z, sample #2-4a (Mo/Cu/Zn/Se) with different Cu ratio ([Cu]/[Zn]

= 1.3) is prepared and is compared with sample #2-4 ([Cu]/[Zn] = 2).

Reaction path for sample #2-4 (Mo/Cu/Zn/Se, [Cu]/[Zn] = 2) at 30–230 C

At the beginning of the measurement for sample #2-4 (Mo/Cu/Zn/Se, [Cu]/[Zn] = 2), only

peak Z is detectable. It can denote Cu, Zn, Cu5Zn8 and CuZn phases in this case. However,

here weak additional peaks for Zn or Cu5Zn8 are not observable, similar to an early stage

of Figure 4.7 for sample #2-3 (Mo/Zn/Cu/Se). Upon comparison with the detectable weak

reflections of Zn in Figure 4.1 (sample #1-1; Mo/Zn/Se)) or Figure 4.12 (sample #2-5;

Mo/Sn/Zn/Se), it is obvious that most of the Zn combines with Cu and forms CuZn during

sputtering, as same as sample #2-3. Thus, peak Z in Figure 4.9 is also mainly due to CuZn

and Cu phases according to the elemental ratio of the precursor. The presence of trace

amounts of Cu5Zn8 and/or Zn in the precursor of sample #2-4 is possible but does not

affect the reaction path due to the Cu diffusion through film, particularly through the Cu–

Zn alloy by dezincification.

At ~190 C, partially crystalline Se suddenly collapses while Se crystallises and melts

Figure 4.9: Time–temperature evolution of powder diffractograms for sample #2-4 (Mo/Cu/Zn/Se) with colour-coded intensities (a.u.) of Bragg reflections. Phases denoted by peaks are as follows: b: ZnSe, l: Cu2Se, m: Cu0.7Zn0.3, Z: Cu, Cu–Zn alloy (CuZn and/or Cu2Zn), and possibly Zn and Cu5Zn8.


33

at ~220 °C. It is shown by sudden weakening of its reflections at ~190 °C in Figure 4.9.

Results for samples #1-1 (Mo/Zn/Se) and #2-5 (Mo/Sn/Zn/Se) which are prepared from

the Zn layer in contact with Se show the absence of diminishing Se reflections before its

melting point. It also indicates that these diminishing reflections at ~190 C are not due to

a growing Zn structure but rather due to Cu beneath the Zn layer. This influence of Cu on

crystalline Se, which is separate from the Se layer, may also be observed in the result for

sample #2-1 (Mo/Cu/Sn/Se). This previous result (see section 4.1.2.1) shows that CuSe

forms at ~200 C before SnSe forms, although the Cu layer is at the bottom of the film.

Moreover, results for samples #2-2 (Mo/Sn/Cu/Se) and #2-3 (Mo/Zn/Cu/Se) prove the

reaction of Cu with crystalline Se. These results – which are the outward diffusion of Cu

and the combination of Cu with crystalline Se, as well as the diminishing Se reflection in

Figure 4.9 – suggest the formation of nanocrystalline CuSe structure which could not be

detected by XRD because of its broad reflection.

While the Se reflections diminish and disappear, peak Z shifts to lower Bragg angles

at 190–230 C and stays at ~43.0° at 230–360 C. Figure 4.10 is the qualitative phase

analysis of peak Z at ~350 °C. According to this phase analysis, the shifting peak Z at

230–360 C implies the co-existence of Cu2Zn (ICDD #00-058-0457) and CuZn phases.

Similar to sample #2-3, dezincification also occurs in the β’-CuZn phase of sample #2-4

intrinsically as the temperature rises, leading to diffusion of decomposed Cu into the film.

Because the Bragg angle of β’-CuZn (peak Z) is the same as that for Cu at ~43.2°, the

metallic Cu may be presented in the film in accordance with its elemental ratio. Therefore,

the sequence of components inside the Cu–Zn alloy under the Se layer may be deduced as

follows: Mo/CuZn/Cu2Zn/Se, Mo/Cu/Cu2Zn/CuZn/Se and so forth. However, the alloy

Figure 4.10: The qualitative phase analysis of peak Z in Figure 4.9 (Mo/Cu/Zn/Se) at ~350 °C. On a basis of Mo reflection at 40.5°, the reflection at around ~43° (the shifting peak Z in Figure 4.9) implies the co-existence of Cu2Zn and CuZn. The red line on the upper part of this Figure indicates the difference between the detected diffraction and the refined diffractions by the references of each phase.


34

composition for peak Z is still uncertain, and the vanishing temperature of this peak Z is

much higher than that for sample #2-3 (Figure 4.7).

To understand better the formation of Cu–Zn alloys (peak Z in Figure 4.9) under the

Se layer before ZnSe formation, another precursor (sample #2-4a) is prepared with the

same sequence of stacked layers but with a different Cu/Zn ratio (1.3 instead of 2). It is

then investigated by in situ analysis, as described in Figure 4.11. The significant

difference between Figures 4.9 (sample #2-4) and 4.11 (sample #2-4a) is the presence of a

curved peak n at 230–350 C in Figure 4.11. Peak m in Figure 4.9 at 360–380 C seems to

be the same reflection as peak n in Figure 4.11, but both have different Bragg angles

(~42.1° for peak m and ~42.2° for peak n). Peak m for sample #2-4 ([Cu]/[Zn] = 2; Figure

4.9) has a Bragg angle lower than that for peak n for sample #2-4a ([Cu]/[Zn] = 1.3;

Figure 4.11).

Reaction path for sample #2-4a (Mo/Cu/Zn/Se, [Cu]/[Zn] = 1.3)

Figure 4.11: Time–temperature evolution of powder diffractograms for sample #2-4a (Mo/Zn/Cu/Se with Cu concentration (Cu/Zn = 1.3) lower than that of sample #2-4 in Figure 4.9) with colour-coded intensities (a.u.) of Bragg reflections. Phases denoted by peaks are as follows: b: ZnSe, h: CuSe, i: l-CuSe2, j: s-CuSe2, l: Cu2Se, n: Cu0.7Zn2, Z: Cu, Cu–Zn alloy (CuZn and/or Cu2Zn), and possibly Zn and Cu5Zn8.

The emerging peak n at ~230 °C and the diminishing peak Z in Figure 4.11 (sample

#2-4a) indicate phase transformation from CuZn (β’ phase) to ε-brass. Here, peak Z is also

mainly produced by the β’-CuZn phase, similar to peak Z in Figure 4.9, because the

additional peaks near 35–39° for Cu5Zn8 and Zn are undetectable in Figure 4.11. Peak n

can denote two kinds of Cu–Zn alloys, since references for these two phases closely

match the lowest Bragg angle of this peak n at ~42.2°. One is for α-brass, (Cu13Zn7)0.2


35

(ICDD #97-062-9457), which is the same as that for Cu65Zn35; the other is for ε-brass,

Cu0.7Zn2 (ICDD #97-010-3153), which can be written as CuZn3 [11]. The absence of peak

n in Figure 4.9 (sample #2-4) suggests that peak n denotes Cu0.7Zn2 instead of Cu65Zn35,

because sample #2-4 can easily compound with Cu65Zn35 owing to the inclusion of

proportions of Cu higher those in than sample #2-4a. According to the study investigated

by H.E. Troiani et al [61], the formation of α-brass from β-brass by dezincification is very

difficult. Instead of the formation of α-brass, the decrease in Zn concentration in β-brass

(Cu-48 at% Zn) leading to a concentration of 38.7 at% Zn (Cu-38.7 at% Zn) is observed

in this study [61]. It indicates the difficulty of formation from β’-CuZn to Cu65Zn35.

Furthermore, one of literature [63] clearly presents that the ε-brass peak at ~42° grows and

shifts to lower angles as the Zn concentration increases in the Cu–Zn alloy, whereas the

peak at ~43° is diminishing, similar to the shift of peak n at at 230–300 °C. Thus, the

denotation of ε-brass for peak n is more plausible than of α-brass. Additionaly, the

decomposed Cu from CuZn in sample #2-4a seems to form not the α-brass but a Cu-rich

CuZn (CurichZn) alloy, which includes Cu concentration lower than α-brass but higher

than β’-CuZn, similar to the formation of Cu-38.7 at% Zn alloy in the study of [61].

The ε-brass phase seems to form at the bottom of film as a result of dezincification

and outward diffusion of Cu. As observed in previous results, Cu diffuses outward to the

film when Cu is mobile. Cu of Cu–Zn alloys can move when it decomposes via

dezincification. Pickering observed the formation of α-brass on the surface of ε-brass

alloy by dezincification when homogenous ε-brass is annealed at 380 C for 5 weeks [62].

In this work, the tendency of Cu to diffuse outward to the film is considerable although

sample #2-4a is not annealed at 380 C for 5 weeks. This tendency may interprete as the

movement of Zn to the bottom of the film by means of outward diffusion of Cu. In other

words, the movement of Zn to the bottom of the film is induced by Cu. In addition,

sample #2-4 suggests the possible formation of nanocrystalline CuSe, which can be

observed by the diminishing Se diffractions at ~190 C before its melting point (see

Figure 4.9); this nanostructure seems to form between metallic Zn and Se layers because

of Cu outward diffusion. Likewise, sample #2-4a also shows a sudden decrease in Se

reflections at ~190 C (see Figure 4.11). While Cu decomposed from the Cu–Zn alloy

diffuses upward and forms nanocrystalline CuSe in between Cu–Zn alloy and Se element,

the remaining (or selectively leached) Zn forcedly moves to the back contact (Mo) due to

the outward diffusion of Cu. For these reasons, the formation of ε-brass at the bottom of

the film in sample #2-4a seems reasonable.

Another notable observation in Figure 4.11 is the shift of peak n depending on the Zn

concentration in ε-brass. After the Se reflections weaken, the β’-CuZn phase (peak Z)

steadily transforms into the ε-brass (peak n) from ~230 C by dezincification and Cu

outwardly diffuses. As soon as Zn decomposes from Cu–Zn alloy during the increase in


36

temperature, peak n shifts to low angles until ZnSe (peak b) appears at ~300 C. After the

emergence of peak b, peak n shifts to high angles again and disappears at ~350 C when

peaks b (ZnSe) reach highest intensities, implying means the consumption of Zn from Cu–

Zn alloy. This result demonstrates the decrease in Bragg angles of peak n with the

increase in Zn concentration of ε-brass.

The last reaction path for sample #2-4a before disappearance of peak n at ~350 C is

shown by the brief appearance of faint peaks h (CuSe) for 340–350 C and the

simultaneous emergence of peak l (Cu2Se) at ~340 C as the Se reacts with decomposed

Cu from Cu–Zn alloys. It intends that the decomposition of ε-brass causes the formation

of Cu2Se at temperature lower than the peritectic decomposition temperature of CuSe.

The observation of the reaction process for sample #2-4a verifies that outward

diffusion of Cu occurs through the Zn layer, forming nanocrystalline CuSe, CurichZn, and

Cu0.7Zn2 phases from upper to bottom parts of Cu–Zn alloy: Mo/Cu0.7Zn2/CurichZn/(nano-

size) CuSe/Se. Because of the low Cu/Zn ratio in sample #2-4a, the Cu0.7Zn2 phase is

formed on the lower side of the Cu–Zn alloy, instead of the formation of Cu2Zn phase on

the upper side of the Cu–Zn alloy.

Reaction path for sample #2-4 (Mo/Cu/Zn/Se, [Cu]/[Zn] = 2) from ~230 C

On a basis of the reaction path for sample #2-4a, the reaction path for sample #2-4

(Figure 4.9) is clarified that the Cu–Zn phases under the Se layer at 230–360 C consist of

Mo/β’-brass/β’+α-brass/α-brass/Se (β’-brass for CuZn and α-brass for Cu2Zn) with an

inverse concentrations between Cu and Zn through Cu–Zn alloy. These are in accordance

with a previous result [62]. There is also little possibility for the Cu0.7Zn2 formation at the

bottom of sample #2-4, but the amount of this phase would be small since the amount of

Cu in sample #2-4 is larger in comparison with that in sample #2-4a. Therefore, only peak

Z denoting CuZn and Cu2Zn is detectable here until ZnSe forms.

At ~360 C, Se reacts with Zn from the Cu–Zn alloy. Because of the use of Zn for the

ZnSe formation, Cu–Zn alloy transforms into another α-brass, Cu0.7Zn0.3 (ICDD #03-065-

9062). As shown in Figure 4.9, peaks b (ZnSe) and m (Cu0.7Zn0.3) appear at the same time

at ~360 C. The simultaneous appearance of peaks b and m signifies the formation of α-

brass (Cu-rich Cu–Zn alloy) with a Zn concentration of ~33–50 at% [9], corresponding to

peak m. This is due to the consumption of Zn from Cu–Zn alloys in accordance with the

elemental ratio of its precursor. Clearly, the phase for peak m is different from that for

peak n because peak m has a Bragg angle lower than that of peak n. In accordance with

the elemental ratios of these two samples, however, peak m includes a higher proportion

of Cu than that in peak n. These two facts signify that peak m does not belong to ε-brass.

Peak m may also denote two kinds of α-brass, one for the (Cu13Zn7)0.2 phase, which is a

kind of a high brass (or yellow brass: Cu65Zn35) [64] and another for a Cu0.7Zn0.3 phase.


37

Although both phases have slightly larger Bragg angle than that for peak m, there is no

exact match to the Cu–Zn alloy. Only two phases have a Bragg angle near this peak m.

Considering the transformation of Cu2Zn into each phase, the chemical reactions can be

written as follows: 7 Cu2Zn → 20 Cu0.7Zn0.3 + Zn and 35 Cu2Zn → Cu65Zn35 + 5 Cu.

According to these two chemical reactions, the Cu0.7Zn0.3 seems to be more reasonable

than Cu65Zn35 because Zn needs to be remained after this transformation of Cu2Zn for the

formation of ZnSe: observing Figure 4.9, peak m also slightly shifts to higher Bragg

angles as peak b strengthens, finally disappearing as soon as peak l (Cu2–xSe) appears,

similar to peak n in Figure 4.11. As shown in Figure 4.11, peak n shifts to low angles as

the proportion of Cu in ε-brass decreases. These similar changes in peaks n and m also

verify the selective leaching of Zn from Cu–Zn alloy for the formation of ZnSe, hence the

correspondence of Cu0.7Zn0.3 with peak m. Meanwhile, Zn from the lower part of the Cu–

Zn alloy layer also reacts with Se by Se diffusion through the film. Therefore, the reaction

path for alloys near ~360 C may be described as following equation (5). Zn has also been

found to have strong tendency to react with Se compared with Cu at high temperature: (v)

Cu–Zn < Zn–Se. The exact temperature for this reaction tendency is determined

depending on the Cu concentration in the Cu–Zn alloy under Se layer, in accordance with

the comparison for the formation temperatures of ZnSe between sample #2-4 and #2-4a.

7 Cu2Zn + Se → ZnSe + 20 Cu0.7Zn0.3 (in the upper side of Cu–Zn alloy) (5)

7 CuZn + Se → 4 ZnSe + 10 Cu0.7Zn0.3 (in the lower side of Cu–Zn alloy)

The left reaction path for sample #2-4 is the formation of Cu2–xSe at ~370 C along

with the disappearance of Cu0.7Zn0.3. Because of the consumption of Cu from Cu–Zn alloy

for the Cu2–xSe formation, peak m (Cu0.7Zn0.3) completely disappears at ~380 C while

peak l (Cu2–xSe) emerges at ~370 C and shifts to the low-angle side, as described in

Figure 4.9. In addition, the shift of peaks l indicates the expansion of the Cu2–xSe structure

into Cu2Se (peaks l).

4.1.2.5 Reactions of Mo/Sn/Zn with Se

Sample #2-5 (Mo/Sn/Zn/Se) replaces the Cu layer of sample #2-4 (Mo/Cu/Zn/Se) with

a Sn layer. This replacement of Cu layer by Sn layer induces a different formation

temperature of ZnSe in the reaction path. In contrast to the previous result showing the

delay of ZnSe formation at ~360 C by a Cu layer in sample #2-4, sample #2-5 presents

the same formation temperature for ZnSe for sample #1-1 (Mo/Zn/Se) at ~290 C. In

other words, the Sn layer beneath the Zn layer does not influence ZnSe formation, in

contrast the Cu layer beneath the Zn layer.


38

Figure 4.12: Time–temperature evolution of powder diffractograms for sample #2-5 (Mo/Sn/Zn/Se) with colour-coded intensities (a.u.) of Bragg reflections. Phases denoted by peaks are as follows: a: Zn, b: ZnSe, c: Sn, d: SnSe, e: SnSe2, E: SnSe and/or SnSe2 (e).

At the beginning of measurement for sample #2-5, Sn (peaks c) and Zn (peaks a) are

obviously detected in the absence of compounds or binary selenides in Figure 4.13. The

undetectable SnSe or ZnSe, which are inconsistent with the phase diagrams of Sn–Se and

Zn–Se alloys, correspond to previous results for samples #1-1 (Mo/Zn/Se) and #1-2

(Mo/Sn/Se). In particular, the growth of the Zn structure along the z-axis is the same as

that for the sample #1-1. This growth may confirm in Figure 4.12 by the shifting of two

peaks at 36.29° for the (002) plane and at 43.23° and for the (101) plane to low Bragg

angles during the increase in temperature.

Meanwhile, the eutectic Sn–Zn alloy gradually forms from ~200 °C in between Sn and

Zn layers by using Sn and part of Zn, in accordance with the Sn–Zn phase diagram [16].

The Sn reflections (peaks c) in Figure 4.12 become weak as soon as the Se reflections

emerge at ~110 °C. The reason of its weakening is unclear and may not reveal in this

study. Subsequently, the Sn reflections vanish at ~200 °C, which is below its melting

point. This temperature for the disappearance of peaks c is in accordance with the eutectic

temperature for Sn–Zn alloy [16]. According to the Sn–Zn phase diagram described in

section 2.1.1.3, eutectic Sn–Zn alloy forms at 198.5 °C when the Zn concentration is

higher than 14.9 at%. Another study also shows the formation of a eutectic Zn–Sn alloy at

around 200 °C by differential scanning calorimetry (DSC) [65, 66]. Consequently, the

disappearance of Sn reflections (peaks c) at ~200 °C in Figure 4.12 clearly indicates the

formation of Sn–Zn alloy as a eutectic mixture. Contrary to the disappearance of Sn

reflections, Zn reflections (peaks a) are still detectable at temperatures higher than


39

~200 °C. These differently detectable reflections of Sn and Zn imply the slow diffusion of

Sn from bottom to top of Zn layer, forming a eutectic alloy. Because of its eutectic

composition (Sn-14.9 at% Zn), it is possible to form this alloy with a small amount of Zn

relative to the amount of Sn. Moreover, Sn and Zn layers are separately deposited in the

precursor of sample #2-5 (Mo/Sn/Zn/Se). That is, when the temperature reaches at

eutectic temperature (198.5 C), the Sn–Zn alloy is formed from most of the Sn and a part

of the Zn with a relatively small amount of Zn in between Sn and Zn layers.

While a part of the Zn on the Sn layer (lower part of Zn layer) forms the eutectic Sn–

Zn alloy, another part of the Zn beneath the Se layer (upper part of Zn layer) seems to

orients toward liquid Se after Se melts at 221 C. Over several measurements of sample

#2-5, three Zn reflections (peak a) vanished gradually at different temperatures. These

different temperatures of disappearance of each reflection are within the range of 290–

320 C. One common observation for these reflections of Zn is the strengthening of one

reflection at ~36° which indicates the (002) plane as soon as Se reflections disappear at

~220 C. Another part of the Zn beneath Se the layer, which is not combined with Sn yet

during the gradual formation of eutectic alloy from the bottom of Zn layer, appears to be

oriented towards the [002] direction, where the melting Se is. The reason of it cannot be

clearly revealed in this study but would be related to the strong affinity of Zn to Se. Upon

comparison of the range of vanishing temperatures of Zn reflections at 260–290 C for

sample #2-6 (Mo/Zn/Sn/Se) in section 4.1.2.6, it seems that the penetration of Sn through

Zn layer from bottom to top for the formation of Sn–Zn alloy seems to have a relatively

slower rate than the diffusion of Sn through Zn from top to bottom of Zn layer. For this

reason, the Zn reflections in Figure 4.12 (sample #2-5) disappear at temperature higher

than that for Figure 4.13 (sample #2-6).

At ~290 C, ZnSe formation starts with a trace amount before Zn reflections disappear

at 290–320 C. Therefore, the emerging peaks b (ZnSe) with faint intensity at ~27° can be

found in Figure 4.12 at ~290 C, while peaks a (Zn) can be clearly detected. This

formation of ZnSe before the disappearance of Zn reflections corresponds well to the

result for sample #1-1 (Mo/Zn/Se). However, temperatures at which Zn reflections vanish

are different (290–320 C for sample #2-5 and ~370 C for sample #1-1). This difference

is caused by the consumption of Zn for the formation of Sn–Zn alloy while a part of the

Zn beneath the Se layer reacts with Se at ~290 C. Zn in sample #2-5 is used for

formation of the eutectic Sn–Zn alloy, whereas Zn in sample #1-1 is used in the steady

growth of its structure before it reacts with Se. Therefore Zn reflections of sample #2-5

vanish at temperatures lower than those for sample #1-1.

At ~330 C, SnSe starts to form gradually, and its reflection becomes observable at

~350 C. Peak d (SnSe) emerging at ~330 C is not clearly seen in Figure 4.12, but when

the diffractions near these temperatures are magnified, the growing peak d with a weak


40

intensity can be observed. This faint peak becomes visible when the temperature reaches

at ~350 C (Figure 4.12). Here, the formation temperature of SnSe at ~330 and ~350 C

can be explained by the liquidus temperature of eutectic Sn–Zn alloy [16]. As described in

section 2.1.1.3, the eutectic Sn–Zn alloy transforms into a liquid phase at around 330 C

(or at higher temperature) when the Zn concentration is near 50 at% (or less). Because of

the elemental ratio of sample #2-5, [Sn]/[Zn] = 1, and the partial consumption of Zn for

ZnSe formation from ~290 °C, the eutectic alloy of this sample becomes liquefied at

around 330–350 C. These liquidus temperatures (330–350 °C) of Sn–Zn alloy are

coincident with the temperatures for the emergence and growth of peak d (SnSe) in Figure

4.12. This signifies that SnSe can form after the eutectic alloy converts into a liquid phase.

In other words, the formation temperature of SnSe depends on the Sn composition in the

Sn–Zn alloy. Thus, the non-uniform distribution of Sn through the eutectic alloy in the

film can also be inferred from the the slowly emerging peak d from ~330 °C; if the Sn

distribution through Sn–Zn alloy is even, the Sn–Zn mixture would melts at the same

temperature and would gradually forms SnSe, as like the beginning of peak E in Figure

4.2 (sample #1-2; Mo/Sn/Se).

Finally, SnSe2 (peak e and E) forms at ~360 C because of the increase in sample

pressure due to evaporation of Se and/or SnSe gas. The intensity of peak d in Figure 4.12

increases at ~360 C as peaks e at 14.43° (not shown) and 44.28° appear, similar to the

results for sample #1-2 (Mo/Sn/Se).

4.1.2.6 Reactions of Mo/Zn/Sn with Se

Sample #2-6 (Mo/Zn/Sn/Se) with inverse sequence of Sn and Zn layers for sample #2-

5 (Mo/Sn/Zn/Se) is prepared. The formation sequences of binary selenides for samples

#2-6 and #2-5 are the same in spite of the different stacking order of elemental layers.

ZnSe formation is earlier than that of SnSe regardless of the sequence of Sn and Zn layers.

The beginning reaction path for sample #2-6 described in Figure 4.13 is exactly the same

as that for sample #2-5 in section 4.1.2.5. Here, Zn (peaks a) and Sn (peaks c) are also

clearly observable at ~30 C, and Sn reflections slightly weaken as soon as Se reflections

appear. Subsequently, Sn reflections disappear at ~200 C for the same reason as that for

sample #2-5; the eutectic Sn–Zn alloy forms at ~200 °C while Zn reflections are still

observable (see section 4.1.2.5 for a more detailed description). Afterwards, Zn reflections

for sample #2-6 show slightly different reaction tendency in comparison with those for

sample #2-5.

After Se melts at ~220 C, the Zn structure grows without any orientation and then

disappears at 260–290 C. In contrast to the strengthening Zn reflection at ~36° in Figure

4.12 (sample #2-5), the three Zn reflections in Figure 4.13 (sample #2-6) do not change


41

Figure 4.13: Time–temperature evolution of powder diffractograms for sample #2-6 (Mo/Zn/Sn/Se) with colour-coded intensities (a.u.) of Bragg reflections. Phases denoted by peaks are as follows: a: Zn, b: ZnSe, c: Sn, d: SnSe, e: SnSe2, E: SnSe and/or SnSe2 (e).

their intensities, similarly to the Zn reflections in Figure 4.1 (sample #1-1; Mo/Zn/Se).

While several measurements were performed to confirm the reaction path for sample #2-6,

the temperature for disappearance of Zn reflections varies with each measurement and

sometimes with each reflection at the range of 260–290 °C. Contrary to the observable Zn

reflections after ZnSe formation in Figures 4.1 (sample #1-1) and 4.12 (sample #2-5), the

Zn reflections in Figure 4.13 completely disappear before ZnSe formation. This signifies

that all of the Zn in sample #2-6 is used to form the eutectic Sn–Zn alloy before the ZnSe

formation. The disappearance of Zn reflections at higher temperature than the vanishing

temperature of Sn reflections in Figure 4.13 also implies the gradual consumption of Zn in

the formation of eutectic Sn–Zn alloy, in contrast to Sn, which is used immediately.

Because of the eutectic Sn–Zn alloy, ZnSe (peak b) is formed at ~350 C. This

formation temperature of ZnSe is higher than that for sample #2-5 (~290 C) probably

because of the liquidus temperature of eutectic Sn–Zn alloy. Contrary to the beginning of

ZnSe formation by metallic Zn (not from Sn–Zn alloy) and Se in sample #2-5, ZnSe is

formed by Zn from eutectic Sn–Zn alloy and Se in sample #2-6 because all of the Zn is

already used in the eutectic alloy formation. This formation of ZnSe in sample #2-6 is

similar to the formation of SnSe in sample #2-5 in terms of the use of element from the

eutectic alloy. In the previous result for SnSe formation, Sn of a eutectic Sn–Zn alloy

reacts with Se after Sn–Zn alloy becomes a liquid phase. In the same manner, Zn from

Sn–Zn alloy also seems to be able to react with Se after the Sn–Zn alloy melts, ie, the

formation temperature of ZnSe is same as the liquidus temperature of eutectic Sn–Zn


42

alloy for sample #2-6. However, the liquidus temperature of this alloy (or the formation

temperature of ZnSe at ~350 C) is slightly higher than the temperature on the Sn–Zn

phase diagram, as indicated by the elemental ratio for sample #2-6 ([Sn]/[Zn] = 1). As

described in section 2.1.1.3, the liquidus temperature is around 330 C when the Sn

composition is 50 at% (Sn-50 at% Zn). Furthermore, the Sn composition in the Sn–Zn

alloy decreases as the liquidus temperature increases. Thus, ZnSe formation at ~350 C

indicates that the Sn-63 at% Zn alloy was formed beneath the liquid Se before the

formation of ZnSe, in accordance with the Sn–Zn phase diagram. It can also say that the

Sn concentration through Sn–Zn alloy is uneven due to the Se diffusion through Zn layer

from top to bottom, as evidenced by the sequence of stacked layers in sample #2-6

(Mo/Zn/Sn/Se) and by its elemental ratio ([Zn]/[Sn] = 1). The movement of Zn from a

Sn–Zn alloy toward to the other element is also observed in other studies [66]. Upon

preparation of the Sn–Zn mixture on the Cu layer, Zn from the Sn–Zn mixture steadily

adheres to Cu and forms a Cu–Zn alloy between the Cu layer and the Sn–Zn mixture [66].

The result from this study [66] confirms the high concentration of Zn on the upper part of

the Sn–Zn alloy due to a Se layer on the Sn–Zn alloy; thus, the higher concentration of Zn

than Sn beneath the Se layer is reasonable. In addition, the liquidus temperature of Sn-63

at% Zn alloy at ~350 C (higher than ~330 C for the Sn-50 at% Zn alloy) is also

reasonable. These results reveal that Zn has stronger tendency to react with Se than Sn as

Zn moves to Se through the Sn–Zn mixture: (vi) Zn–Sn < Zn–Se.

Lastly, SnSe (peaks d and E) forms at ~360 C. After a rise of ~10 C to ~370 C,

SnSe2 (peaks e and E) is formed and becomes a main phase for the Sn–Se alloy. This

reaction path is clearly described in Figure 4.13 by peaks d, e and E: After the appearance

of peak d and E at ~360 °C, peak d disappears at ~380 C while peak e appears at ~370 C

and grows. The sudden transformation of SnSe to SnSe2 signifies that the sample pressure

at this temperature is high enough to form SnSe2, in accordance with the result for sample

#1-2 (Mo/Sn/Se). This previous result for sample #1-2 suggests the influence of sample

pressure on a SnSe2 formation. Becasue the melting point of Se is at ~220 C and the first

reaction temperature of Se with Zn is at ~350 C, some of the Se evaporates, increasing

the sample pressure as SnSe2 forms. Although the Sn layer is deposited beneath the Se

layer in the precursor (Mo/Zn/Sn/Se), it does not react with Se at ~230 C, in contrast to

the result for sample #1-2 (Mo/Sn/Se). When the eutectic Sn–Zn alloy melts so that Zn

reacts with Se at ~350 °C, Sn does not react with Se at the same time with Zn. On the

basis of these data, the inferred reaction tendency of Sn is follows: (vii) Sn–Se < Zn–Sn.

4.1.3 Results of experiments on single and double metal layers

The results for the Se reaction with one and two metallic elements demonstrate the


43

dependence of processes for binary selenide formation on the sequence of stacked layers

in the precursor. Formation temperatures are summarised in Table 4.1. Several

characteristics of reactions can be derived from the different reaction paths of each sample,

as described in the sections below.

Table 4.1: The formation temperatures of each selenide measured in one-/two-metal samples. The unit for the temperature number is the degree Celsius [°C]. Depending on the sequence of elemental stacking layers in the precursor, the formation temperatures of ZnSe and SnSe are changed.

One metallic

layer

Two metallic layers

Cu–Sn–Se Cu–Zn–Se Sn–Zn–Se

Sample

#

#1-1 #1-2 #2-1 #2-2 #2-3 #2-4 #2-4a #2-5 #2-6

Zn/Se Sn/Se Cu/Sn/Se Sn/Cu/Se Zn/Cu/Se Cu/Zn/Se Sn/Zn/Se Zn/Sn/Se

CuSe ~200 ~30 ~30

l-CuSe2/

s-CuSe2

~220/

~230

~190/

~220

~170/

~210

Cu2Se 380–400 380–400 ~380 ~370 ~340

ZnSe ~290 ~290 ~360 ~300 ~290 ~350

SnSe ~230 ~250 ~230 ~330 ~360

CTSe ~290 ~290

4.1.3.1 Influence of pressure on Sn–Se alloy formation

The formation of Sn–Se alloy is significantly influenced by the sample pressure, in

accordance with the result of sample #1-2 in section 4.1.1.2. This previous result presents

the change in Sn–Se alloy depending on the sample pressure. Because the melting point of

Se is 221 °C and because SnSe can be evaporated during annealing [5, 53], sample

pressure steadily increases until the clamped sample ruptures. This sample consists of two

pieces from one sample in a face-to-face state set in a sample holder (see section 3.2.1)

and then applies heat. This sample rupture changes the sample pressure during the

measurement. At this point, transformation of SnSe2 to SnSe in the sample #1-2 is clearly

observed. Thus, the sample pressure needs to be high enough to form SnSe2 in the film.

4.1.3.2 Reactive Cu

Cu is sufficiently reactive to allow compounding of CuSe in the precursor, in contrast

with the Zn and Sn which do not form ZnSe and SnSe by evaporating Se layer during the

preparation of each sample. Although ZnSe is the alloy phase that most easily forms

among the binary selenides in accordance with the Ellingham diagram [22] and the values

of each free energy [23], ZnSe is undetectable at room temperature in sample #1-1 or #2-5,


44

whereas CuSe is detectable at room temperature in sample #2-2 and #2-3 (see Table 4.1).

When the Cu layer is sputtered on the other metallic layer or vice versa during the

preparation of samples, Cu easily reacts with other metallic elements, too, and forms a

Cu–metal alloy in the precursor, such as Cu6Sn5 in samples #2-1 and #2-2, or CuZn in

samples #2-3 and #2-4. The formation of alloys based on Cu in the precursor signifies that

Cu is a highly reactive element among the four components of CZTSe.

4.1.3.3 Outward diffusion of Cu: blocked only by a Zn layer

The tendency of Cu for outward diffusion may be observed in our data, but Cu

diffusion is interrupted only by a Zn layer. As discussed in other works [61, 62, 67],

outward diffusion of Cu is easily observed. In particular, Pickering clearly shows the

outward diffusion of Cu in the formation of α-brass (Cu-rich Cu–Zn alloy) from γ-/ε-brass

(Cu-poor Cu–Zn alloy) [62]. Similar to these results, Cu diffused through the Sn layer into

the Se layer, forming a Cu3Sn phase in sample #2-1 (Mo/Cu/Sn/Se). Consequently, Cu

could react with Se earlier than Sn reacts. In contrast to this result, Cu could not diffuse

through Zn layer into the Se layer in sample #2-4 (Mo/Cu/Zn/Se); thus Cu–Se alloy

cannot form before ZnSe formation. Although Cu could diffuse through the Zn layer by

forming a Cu2Zn phase under the Se layer, it could not react with a liquid Se which is

located on the Cu–Zn alloy (Mo/CuZn/Cu2Zn/liquid-Se) up to ~360 C, as shown in Table

4.1. That is, outward diffusion of Cu is blocked by Zn.

In the case of the sulfur (S), which is chemically similar to Se, outward diffusion of

Cu is also observed. According to an investigation by Buckel et al. [68], pieces of Cu

metal and S vapour react at 470 C, forming an inner hole and an outer CuSe mantle in the

pieces of Cu together with an intermediate zone of Cu1.8S [69]. Similarly, the reaction

sequence of Cu–Se alloy also features outward diffusion of Cu by transformation of l-

CuSe2 into s-CuSe2. This reaction sequence of Cu–Se alloy is discussed in section 4.1.3.5.

4.1.3.4 Induced movement of Zn to the back electrode by Cu

Zn diffusion to the inner layer is only observed in the Cu–Zn alloy because of

dezincification and outward diffusion of Cu. This diffusion indicates that in prepared

samples with a Cu layer in contact with Zn layer, Zn moves to the bottom of the film as

follows: i) Cu–Zn alloy, mainly the CuZn phase and a trace of Cu5Zn8, forms in the

precursor, while Cu and Zn layers were deposited on the substrate. ii) Upon heating of the

sample, Zn atoms selectively leach from the Cu–Zn alloy via dezincification. iii) During

dezincification, the dealloyed Cu diffuses into the outward layer of the Cu–Zn alloy. iv)

Due to the Cu diffusion, Zn is pushed out from its position to the inward layer of the Cu–


45

Zn alloy. The co-existence of Cu2Zn and CuZn (Figure 4.10) in sample #2-4 at 230–

360 °C confirms this subsidence of Zn, although the ratio of Cu/Zn in sample #2-4 is

equal to 2. Moreover, sample #2-4a, which has a lower ratio of Cu/Zn ([Cu]/[Zn] = 1.3)

than of sample #2-4, presents the transformation of CuZn into ε-brass (such as Cu0.7Zn2)

at the lower part of Cu–Zn alloy explained in section 4.1.2.4. Pickering also proves that

Cu and Zn interdiffuse in the Cu–Zn alloy via transformation of ε-brass into α-brass on

the outer layer of Cu–Zn alloy (ε-brass) [62]. Our results and those in the literature

confirm this subsidence of Zn along with the outward diffusion of Cu: in other words, the

interdiffusion of Cu and Zn through Cu–Zn alloy.

4.1.3.5 Reaction sequence of Cu–Se alloys

This section describes the reaction sequence of Cu–Se alloys in detail. The formation

reaction of Cu–Se alloys in the results of this study presents always the same pattern

depending on the temperature. This is in accordance with the Cu–Se phase diagram when

the Se concentration is near 50 at% [18]. This process may be observed in samples #2-2

and #2-3, which have a Cu layer in contact with Se, such as Mo/[metal]/Cu/Se. Although

the Cu layer in sample #2-1 (Mo/Cu/Sn/Se) is deposited on the bottom of the film, this

Cu–Se process is also observed because of the diffusion of Cu simply through the Sn layer.

According to these results for three samples, the transformation of Cu–Se alloys follows

reaction (6):

CuSe + Se(l) → l-CuSe2 → s-CuSe2 → CuSe + Se → Cu2–xSe + Se(g) (6)

In fact, the CuSe2 phase is not divided in the Cu–Se phase diagram [18], but two

different kinds of CuSe2 reflections are clearly detected in our results in Figures 4.5

(sample #2-2) and 4.7 (sample #2-3). One reflection (peak i, l-CuSe2) matches closely

ICDD #97-002-5717, which represents a relatively large unit cell of this structure, and

another (peak j, s-CuSe2) matches ICDD #00-019-0400, which has a relatively small unit

cell. Thus, the change in peak i (l-CuSe2) to peak j (s-CuSe2) observed in Figure 4.7

indicates a decrease in size of the CuSe2 structure. As mentioned in section 4.1.3.2 and

Table 4.1, CuSe is detectable at room temperature in Mo/[metal]/Cu/Se samples because

of the reactive element of Cu. Therefore, during the measurement, Cu steadily diffuses

through the Se layer and alloys with crystallised Se, forming CuSe in abundance and l-

CuSe2. When the temperature reaches ~220 C, which is near the melting point of Se

(221 °C), Se diffuses into the inner Cu–Se alloy, resulting in transformation of the l-CuSe2

into s-CuSe2.

This decrease in size of CuSe2 structure seems to be caused by Se diffusion into the


46

film. When Cu–Se alloy includes certain amount of Se, the Se starts to diffuse through

film, transforming the l-CuSe2 into the s-CuSe2. This Se diffusion from CuSe2 can be

confirmed by comparison of formation temperatures between s-CuSe2 and other binary

selenides, described in Table 4.1. In the case of sample #2-2 (Mo/Sn/Cu/Se), SnSe forms

at ~230 C after s-CuSe2 forms at ~220 C. Likewise, ZnSe forms at ~290 C after s-

CuSe2 is compounded at ~210 C in sample #2-3 (Mo/Zn/Cu/Se). It signifies that the

metallic element of Mo/[metal]/Cu/Se sample can be selenised by Se diffusing from the

Cu–Se alloy after the formation of s-CuSe2.

Afterwards, the s-CuSe2 phase decomposes into CuSe at ~330 C, which is near the

peritectic decomposition temperature of CuSe2 (332 C) [18]. At this point, the CuSe2

reflections weaken before its melting point because of the increase in Cu concentration (or

because of the decrease in Se concentration) in the Cu–Se alloys (eg, Figures 4.4 for

sample #2-1 and 4.7 for sample #2-3); the reflections of CuSe2 do not disappear suddenly

but decrease their intensity and then vanish.

CuSe decomposed from CuSe2 decomposes again at ~380 C into a Cu2Se (or Cu2–xSe)

phase. This formation temperature corresponds to the peritectic decomposition temper-

ature of CuSe at 379.3 °C [18]. This β-Cu2Se (or β-Cu2–xSe) appears sometimes at

~370 C (eg, sample #2-4 (Mo/Cu/Zn/Se)), which is lower than its decomposition

temperature, only when it forms from the decomposition of Cu–Zn alloy. In the case of

sample #2-1 and #2-2, this Cu2Se is detected only in the range of ~20 C (380–400 °C)

after its formation at ~380 °C because of the consumption of this phase in CTSe formation.

In this section, the notable reaction process is Se diffusion into the film after the

formation of s-CuSe2, resulting in formation of another binary selenide such as ZnSe or

SnSe after s-CuSe2 formation. In other words, s-CuSe2 does not disturb but facilitates the

reaction between Se and other metallic element which is under the Cu layer.

4.1.3.6 High affinity of Se to Cu

Cu–Se alloy, especially the s-CuSe2 phase, facilitates Se diffusion into the film,

leading to the reaction of Se with other metallic element. When the Cu layer is deposited

between metallic and Se layers, as in the sequence Mo/[metal]/Cu/Se, Cu does not disturb

the formation of metal selenide, but rather enhances it as if there is no Cu element in the

sample. Details in Table 4.1 show that the formation temperature of SnSe in sample #2-2

(Mo/Sn/Cu/Se) is same as that for sample #1-2 (Mo/Sn/Se) which presents no obstacle for

the reaction of Se with Sn. Contrary to this, the formation temperature of SnSe for sample

#2-5 (Mo/Sn/Zn/Se), which replaces the Cu layer of sample #2-2 with a Zn layer,

increases. ZnSe formation also has a tendency similar to SnSe formation. ZnSe in sample

#2-3 (Mo/Zn/Cu/Se) also has the same formation temperature as that for sample #1-1


47

(Mo/Zn/Se) as if the Cu layer does not exist in sample #2-3 for the ZnSe formation. In

contrast, ZnSe formation in sample #2-6 (Mo/Zn/Sn/Se) requires a temperature higher

than that for sample #1-1. On the basis of the formation temperatures of ZnSe and SnSe,

Se diffusion seems to be promoted through Cu layer by the formation of Cu–Se alloy,

especially the s-CuSe2 phase as mentioned in section 4.1.3.5.

4.1.3.7 Delayed ZnSe formation by the Cu layer beneath the Zn layer

As described in section 4.1.3.3, the Zn layer interrupts the outward diffusion of Cu,

thereby delaying ZnSe formation. As shown in Table 4.1, samples #2-3 (Mo/Zn/Cu/Se)

and #2-5 (Mo/Sn/Zn/Se) can produce ZnSe at ~290 C. The reason for ZnSe formation at

~290 C in sample #2-3 is the easy diffusion of Se through Cu, as described in section

4.1.3.6. The case of sample #2-5 seems reasonable because Zn layer is deposited on the Se

layer on the upper part of its precursor. In contrast to sample #2-5, sample #2-4

(Mo/Cu/Zn/Se) shows a formation temperature for ZnSe higher than that for sample #2-5,

although the Zn layer is in contact with the Se layers in the precursor.

On the basis of the reaction path for sample #2-4, it seems that ZnSe formation is

delayed by the high concentration of Cu on the upper side of the Cu–Zn alloy (see section

4.1.2.4). More likely the Cu in Cu–Zn alloy interrupts the reaction between Zn and Se. As

described in section 4.1.3.4, Cu and Zn interdiffuse in the Cu–Zn alloy below the Se layer,

causing the Cu distribution to decline from the top to the bottom of the Cu–Zn alloy.

Therefore, Se is in contact with a relatively large amount of Cu and a relatively small

amount of Zn before it reacts with Zn from Cu–Zn alloy at ~360 °C. It seems that Cu

disturbs the reaction between Se and Zn in some way. Additionally, sample #2-4a

(Mo/Cu/Zn/Se), which has a proportion of Cu in the precursor ([Cu]/[Zn] = 1.3) lower

than that in sample #2-4 ([Cu]/[Zn] = 2), forms ZnSe at ~300 °C at temperature lower

than that for sample #2-4 (see Table 4.1). The comparison in formation temperatures of

ZnSe between sample #2-4 and #2-4a clearly indicates that the higher concentration of Cu

in the Cu–Zn alloy under Se layer leads the higher temperature of ZnSe formation. Thus,

the formation temperature of ZnSe decreases as the Cu concentration in the Cu–Zn alloy

decreases. In the case of sample #2-3 (Mo/Zn/Cu/Se), because of the consumption of

around half of the Cu for the formation of Cu–Se alloy before ZnSe formation, ZnSe

could form at ~290 C despite that the proportion of Cu in the initial precursor is the same

as that for sample #2-4. Thus, the formation of ZnSe may be delayed by increasing the Cu

concentration in the Cu–Zn alloy under Se layer.

Another reason for the delay of ZnSe formation is the formation of eutectic Sn–Zn

alloy beneath the Se layer, in accordance with the result for sample #2-6 (Mo/Zn/Sn/Se).

In the case of sample #2-6, ZnSe is formed at ~350 C (see Table 4.1) because Zn is


48

trapped in the eutectic Sn–Zn alloy. Details of the reaction path for this sample (see

section 4.1.2.6) indicate that the surface of this alloy near the Se layer has a Zn

concentration slightly higher than that in other areas because of the slow movement of Zn

to Se in the eutectic Sn–Zn mixture. Consequently, the liquidus temperature of Sn–Zn

alloy increases and that causes the reaction of Se with Zn at ~350 °C, at temperature

higher than that for sample #2-5 (~290 °C). This reaction also verifies that ZnSe may form

after the Sn–Zn alloy becomes a liquid phase. It implies that the Zn concentration of the

Sn–Zn alloy also can influence on ZnSe formation.

From these results, the concentration of other metallic elements on the upper side of

metal–Zn alloy (or under Se layer) influences the formation temperature of ZnSe. To

delay the formation of ZnSe, the following sequences of layers are recommended: a Cu

layer beneath Zn/Se layers (Cu/Zn/Se) or a Sn layer on the Zn layer (Zn/Sn/Se). However,

a certain amount of ZnSe in the sequence Cu/Zn/Se can remain at the bottom of the film

because Zn forcedly moves to the back contact via dezincification of Cu–Zn alloy and

outward diffusion of Cu (see section 4.1.3.5).

4.1.3.8 Delayed SnSe formation by Cu and Zn contents of alloy

The formation of Sn–Se alloy is easily influenced by neighbouring metallic elements.

Only when the Cu layer is deposited between Sn and Se layers, as in sample #2-2

(Mo/Sn/Cu/Se), SnSe can form at ~230 C, similar to the result for sample #1-2

(Mo/Sn/Se) which has no element interrupting SnSe formation. The reason for the

identical formation temperature between samples #2-2 and #1-2 is the easy diffusion of Se

through Cu (see section 4.1.3.6). Except sample #2-2, all two-metal samples show a

formation temperature of SnSe higher than that for sample #1-2, as shown in Table 4.1.

SnSe formation in sample #2-1 (Mo/Cu/Sn/Se) occurs at ~250 C, after the disappearance

of Cu3Sn (see section 4.1.2.1). In the case of sample #2-5 (Mo/Sn/Zn/Se), SnSe also forms

at ~330 C when the eutectic Sn–Zn alloy becomes a liquid phase (see section 4.1.2.5).

Both samples show SnSe formation after decomposition of Sn from a metal–Sn alloy.

Contrary to these two samples, SnSe in sample #2-6 (Mo/Zn/Sn/Se) forms at ~360 C

after ZnSe forms at ~350 C. This sequence of formation between ZnSe and SnSe phases

is in accordance with the free energy of these two phases. Because the free energy of SnSe

is higher than that of ZnSe [22-23], the ZnSe structure can form more easily than SnSe

can. Therefore, the appearance of SnSe after ZnSe formation is reasonable. Moreover, the

free energy of SnSe is higher than that CuSe. It means that SnSe is the most difficult alloy

to compound among the binary selenides. Therefore, SnSe forms after the formation of

ZnSe and Cu–Se alloys. In other words, the formation temperature of SnSe is significantly

influenced by neighbouring metallic layers.


49

These two cases reveal that SnSe cannot form unless Sn decomposes from the metal–

Sn alloy and unless other binary selenides (CuSe and ZnSe) form. In any case,

combination of Sn with Se seems unfavourable.

4.1.3.9 Crystalline CuxSey phase determines the rate of Cu2SnSe3

formation

When CuSe2 decomposes into CuSe in the Cu–Sn–Se samples (#2-1 and #2-2), the

formation rate of CTSe varies with the Cu–Se phase. The CTSe phase forms at ~290 C

regardless of the sequence of Cu and Sn layers, but its formation rate is very low, as

shown by a weak reflection of CTSe (peak k) in Figure 4.4 and 4.5 (see sections 4.1.2.1

and 4.1.2.2). During detection of CuSe2 reflections, the intensity of the CTSe reflection is

weak and slowly grows. After CuSe2 decomposes into CuSe and Se at ~330 C, the

formation rate of CTSe increases, as confirmed by the rapidly increasing intensity of

CTSe reflection. This indicates that CTSe usually forms from CuSe and SnSe, instead of

CuSe2 and SnSe. In other words, when CuSe2 comprises the CTSe phase, CTSe forms

slowly in comparison with the consumption of CuSe. Therefore, these reactions can be

described by equations (2) and (3) as described in section 4.1.2.1.

Similarly, the rate of CTSe formation may vary with the consumption of SnSe and the

Sn liquid phase. Because the results of the Cu–Sn–Se samples show only the consumption

of a SnSe phase, however, this study cannot confirm the influence of liquid Sn and a SnSe

phase on the rate of CTSe formation.

4.1.3.10 Conclusion: The tendency of four elements to react with each

other

Two formulas (I) and (II) for the tendency of the reaction between four elements may

be derived from the formulas (i)-(vii) which are revealed during the observation of the

reaction path for two-metal samples over section 4.1.2:

Sn–Se < Cu–Sn < Cu–Se < Cu–Zn < Zn–Se (I)

Sn–Se < Zn–Sn < Zn–Se (II)

The strong tendency of Zn to react with Se than with Cu (Cu–Zn < Zn–Se) is

applicable in the temperature range of 290–360 °C, depending on the Cu concentration in

the Cu–Zn alloy under Se layer. As the Cu–Zn alloy near Se has higher concentration of

Cu, the applicable temperature for this reaction tendency is getting higher. It seems that

the Cu in Cu–Zn alloy disrubs the reaction between Zn and Se as much as the amount of


50

Cu in the Cu–Zn alloy under Se. In particular, when the proportion of Cu to Zn in Cu–Zn

alloy near the Se is double (Cu/Zn = 2; Cu2Zn), the applicable temperature for this

reaction tendency is above ~360 °C, in accordance with the result for sample #2-4

(Mo/Cu/Zn/Se). These two formulas (I) and (II) for the reaction tendency may be

expressed as formula (III) because the Zn–Sn alloy is formed not as a crystallised

structure, but as a eutectic mixture.

Sn–Se < Zn–Sn < Cu–Sn < Cu–Se < Cu–Zn < Zn–Se (III)

In fact, the strong tendency of Sn to react with Zn than with Cu (Zn–Sn < Cu–Sn) is

not confirmed in this study. It is difficult to verify because of the strong affinity of Cu to

Zn: according to the Cu–Zn–Sn phase diagram [17], only Cu–Zn and Cu–Sn alloys are

detectable, whereas the Zn–Sn mixture does not appear in this phase diagram.

Furthermore, when the Sn–Zn mixture is bonded to the Cu layer in a sample and when

heat is supplied to this sample, a Cu–Zn alloy and a separated pure Sn metal are observed

[66]. Therefore the reaction tendency of Sn to metallic elements (Zn–Sn < Cu–Sn) cannot

be identified.

4.2 Investigation of ternary metal systems with Se

51


4.2.1 Correlation of delayed ZnSe crystallisation with the reaction sequence of selenides

According to the Ellingham diagram for selenides [22] or free energies of binary

selenides [23], Zn is more reactive in selenisation than are the other metallic components,

Cu and Sn, regardless of temperature. This signifies that ZnSe can be easily formed in the

sample when three metallic components are co-deposited and are annealed in a Se

atmosphere. That is, once ZnSe alone crystallises well, covalently bonded ZnSe does not

easily react with other phases formed during the heating process inside the film at

moderate temperatures. On the basis of this theory, the correlation between the sequence

of formation of components for CZTSe formation and the residual amount of ZnSe in the

synthesised film is observed with one hypothesis in this section. If the chalcogen Se reacts

with Zn before other components of CZTSe form a compound, then a certain part of ZnSe

alloy develops into a stable structure and subsequently does not easily combine with other

phases into kesterite, causing residual ZnSe in the kesterite film.

To test this hypothesis, three samples with different contact areas between Zn and Se

layers are prepared as follows: Mo/Cu/Sn/Zn/Se for sample #3-1, Mo/Zn/Sn/Se/Cu for

sample #3-2, and Mo/Zn/Sn/Cu/Se for sample #3-3. Because the Sn layer is typically

deposited as a discontinuous island-like structure [70, 71] instead of a closed metal layer,

Zn is in local contact with Cu in sample #3-1 after preparation of the element stacks, with

Se in sample #3-2 and with Cu in sample #3-3. Therefore, Zn in sample #3-2 partially

separates from the Se layer because of the Sn layer, whereas a maximum amount of Zn is

in contact with Se in sample #3-1. Thus, sample #3-2 is intermediate between samples #3-

1 and #3-3 in terms of the contact areas between Zn and Se layers.

After observation of each reaction by in situ analysis in sections 4.2.1.1–4.2.1.3, the

remaining amount of ZnSe in each sample is confirmed by Raman analysis in section

4.2.1.4. Through comparison of the results, the hypothesis is confirmed in section 4.2.1.5.

4.2.1.1 Reactions of Mo/Cu/Sn/Zn/Se

The initial precursor of sample #3-1 (Mo/Cu/Sn/Zn/Se) is composed of Sn, CuZn,

possibly Cu, and a trace amount of Cu5Zn8 without any Cu–Sn alloys although a Sn layer

is deposited on the Cu layer. The in situ XRD diffractogram, which clearly presents peaks

c (Sn) and peak Z at ~30 C, verifies some of these components, as described in Figure

4.14. The single peak denoted Z indicates not only elemental Cu, but also the presence of

Zn, Cu5Zn8, CuZn and Cu6Sn5 due to several overlapping peaks of the strongest Bragg


52

reflections of these phases at ~43.2°. Some of these phases can, in principle, be

distinguished from the others by additional weak reflections at angles of the Zn phase at

36.29° and 39.00° (ICDD #97-005-2543), the Cu5Zn8 phase at 34.98° and 37.89° (ICDD

#97-000-2092), and the Cu6Sn5 phase at 30.15° and 35.20° (peaks b). From these

references for the phases, this precursor is precisely identified by ex situ XRD analysis.

The analysis shows only a trace of Cu5Zn8 phase without any expected Zn reflections, as

shown in Figure 4.15. This analysis implies that most of the Zn had combined with Cu

through an island-like structured Sn layer upon Zn sputtering, similar to the precursor of

sample #2-3 (Mo/Zn/Cu/Se) described in Figure 4.8. The observable, faint reflection near

~35.3° in Figure 4.15 seems to be a distinguishable reflection for Cu6Sn5. However

another distinguishable reflection of Cu6Sn5 at 30.15°, which has intensity higher than that

for the peak at 35.20°, is not detected in Figure 4.14. Considering the clrearly detectable

peak f at 30.15° (Cu6Sn5) in Figure 4.4 (sample #2-1, Mo/Cu/Sn/Se), Cu6Sn5 is obvously

not compounded in this precursor. Here, the absence of a Cu–Sn alloy in sample #3-1 at

~30.1 C is rationalised by the stronger affinity of Cu to Zn than to Sn, as described in

section 4.1.3.10: Cu–Sn < Cu–Zn.

Upon heating of sample #3-1, Sn combines with Cu into Cu6Sn5 at ~190 C, and all of

the Cu–Zn alloys in the film transforms into CuZn. Figure 4.14 at 30–220 C indicates

that as soon as Se reflections appear by its crystallisation, Bragg reflections of the metallic

Sn phase (peaks c) gradually diminish. Afterwards, peaks c (Sn) suddenly disappear

simultaneous with the appearance of peak f (Cu6Sn5) at ~30.1° when the temperature

Figure 4.14: Time–temperature evolution of powder diffractograms for sample #3-1 (Mo/Cu/Sn/Zn/Se) with colour-coded intensities (a.u.) of Bragg reflections. Phases denoted by peaks a-Z are as follows: c: Sn, d: SnSe, f: Cu6Sn5, h: CuSe, K: ZnSe and Cu2ZnSnSe4 (o), o: Cu2ZnSnSe4, Z: Cu, Cu5Zn8, CuZn, Cu6Sn5 (f), and/or α-brass (Cu2Zn).


53

Figure 4.15: GIXD diffractogram of sample #3-1 (Mo/Cu/Sn/Zn/Se) before heat treatment. This analysis is performed in the atmosphere at room temperature. A trace of Cu5Zn8 in the precursor is observed without any trace of Zn [72].

reaches ~190 C upon formation of Cu6Sn5 by Sn and Cu. This formation temperature of

Cu6Sn5 at ~190 °C seems to correlate with the temperature of Cu6Sn5 transformation from

the η’-Cu6Sn5 phase to the η-Cu6Sn5 phase at 186–189 °C [13]. (see section 2.1.1.2). At

optimised temperature for the formation of a hexagonal η-Cu6Sn5 pahse, the separated

metallic Sn seems to be able to combine with Cu, thus the appearance of peaks f (Cu6Sn5)

at ~190 °C in Figure 4.14. The reaction path of Cu – Zn alloys near this temperature can

be inferred from the Sn–Zn–Cu phase diagram [17] which is revealed by Chou et al. This

study [17] investigates the equilibrium phases in the Sn–Zn–Cu phase at 210, 230, and

250 °C. According to this study, when any Cu–Sn alloy is detected in the Sn–Zn–Cu alloy

together with the Cu–Zn alloy, this Cu–Zn alloy is always observed as a β’-CuZn phase

regardless of Zn concentration [17]. It indicates that other Cu–Zn alloys, such as Cu5Zn8

or Cu2Zn, in the Sn–Zn–Cu phase do not co-exist with the Cu–Sn alloys but convert into

the CuZn as the Cu and Sn combine. This result confirm that all kinds of Cu–Zn alloys in

sample #3-1 transform into the CuZn phase at this temperature because of the appearance

of Cu6Sn5. After these transitions, Se melts at 221 C along with the disappearance of Se

reflections at ~220 C (Figure 4.14).

At ~330 C, ZnSe forms by melting of Se and dealloying of Zn from CuZn by

dezincification. ZnSe formation is confirmed by the appearance of peak K at ~330 C. In

fact, this peak K may indicate not only ZnSe but also CTSe and CZTSe because of the

overlap of Bragg reflections of these phases at ~27°. Only CZTSe can be distinguished by

two weak reflections at 35.28° and 36.16° (ICDD #97-009-5117). Although ZnSe and

CTSe cannot be clearly identified from the XRD diffractogram, the beginning of peak K

in Figure 4.14, which is denoted by a ‘K(b)’, clearly corresponds to ZnSe for two reasons.

One reason is the unchanging intensity of peak f (Cu6Sn5) while peak K emerges. If this

peak is produced by the CTSe phase, then the components of CTSe (at least Sn) derive

from Cu6Sn5 by decomposition of Cu6Sn5. Consequently, the intensity of peaks f (Cu6Sn5)


54

would gradually decrease as peak K strengthens. However, the intensity does not change

until ~350 C, ie, the beginning of peak K does not belong to the CTSe phase. Another

reason is the dealloying of Zn from CuZn by dezincification and the strong affinity of Se

to Zn. As the temperature rises, Zn selectively leaches from β’-CuZn by dezincification.

Because Se has stronger tendency to combine with Zn than with Cu [22, 23] and because

the proportion of Cu to Zn in the Cu–Zn alloy is less than double (see section 4.1.3.7), Se

melted at 221 °C may react with the dealloyed Zn at ~330 °C. For these two reasons, the

beginning of peak K obviously indicates ZnSe (peak b).

While the Zn dealloyed from β’-CuZn reacts with Se at ~330 °C, the remaining Cu

adheres not to Cu6Sn5 but to CuZn, in agreement with the reaction tendency of Cu to

metallic elements (Cu–Sn < Cu–Zn), as described in section 4.1.3.10. Thus, a Cu-rich Cu–

Sn alloy such as Cu41Sn11 is not observed in Figure 4.14. However, the reflection for Cu-

rich Cu–Zn alloy, such as α-brass, is also not detected at this temperature. This

undetectable reflection of the Cu-rich Cu–Zn alloy is consistent with the result for sample

#2-4a, which has a Cu/Zn ratio of 1.3 (see section 4.1.2.4). This previous result shows the

disappearance of peak Z (CuZn) along with the appearance of peak n (ε-brass) in Figure

4.11 (sample #2-4a). Although CurichZn forms on the upper part of Cu–Zn alloy in sample

#2-4a during the formation of ε-brass on the lower part, this CurichZn phase could not be

found in Figure 4.11. This CurichZn is the Cu-rich Cu–Zn alloy which has has greater

proportion of Cu relative to that of β’-CuZn and less than that of Cu2Zn (CuZn < CurichZn

< Cu2Zn). In contrast, peak Z remains in Figure 4.9 (sample #2-4, [Cu]/[Zn] = 2),

denoting the co-existence of CuZn and α-brass (Cu2Zn). These two results intend that the

reflection of CurichZn is undetectable whereas the reflection of α-brass is detectable in the

diffractogram. On the basis of these previous results, the undetectable α-brass in Figure

4.14 seems to imply the formation of CurichZn, similar to the result for sample #2-4a.

Another plausible interpretation is the overlapped reflection of α-brass on the peak f

(Cu6Sn5) at ~43°.

One peculiar observation with sample #3-1 and sample #2-4a is the undetectable ε-

brass (peak n) in Figure 4.14 (sample #3-1) in contrast to the observable peak n in Figure

4.11 (sample #2-4a). This difference implies the formation of eutectic Sn–Zn alloy on the

lower part of the film in sample #3-1 until the ZnSe forms at ~330 °C. Considering the

ratio of metallic elements (Cu:Zn:Sn = 1.8:1.2:1) and alloys formed after Cu6Sn5

formation in sample #3-1, as well as the tendency of the reaction (Cu–Sn < Cu–Zn), Sn

should remain in the sample through a reaction described by equation (7) – this equation

considers only the proportion of compounding phases at ~190 °C due to the presence of

the trace amount of Cu5Zn8 in the initial precursor:

18 Cu + 12 Zn + 10 Sn → 12 CuZn + Cu6Sn5 + 5 Sn (7)


55

Because of this remaining Sn, the eutectic Sn–Zn mixture can form in the film before the

ZnSe formation. Therefore, neither α-brass nor ε-brass is clearly detected while Cu6Sn5

(peak f) is obviously observed in Figure 4.14.

At ~350 C, CuSe is formed by Cu dealloyed from Cu6Sn5 and Cu–Zn phases.

Simultaneously, SnSe forms at relatively low rate in comparision with the formation rate

of CuSe. This reaction pathway is observed from the rapidly weakening peaks f (Cu6Sn5)

at ~350 C as soon as the appearance of peaks d (SnSe) and h (CuSe). Peak d steadily

strengthens, starting at intensity lower than that of peak h. The reason for this slow growth

of SnSe reflection (peak d) seems to be the strong affinity of Se to Cu than to Sn. When

Cu6Sn5 decomposes, Se prefers to react with Cu than Sn, in accordance with the reaction

tendency (Sn–Se < Cu–Se; see section 4.1.3.10). For this reason, the formation rate of

SnSe is slower than the formation rate of CuSe at this temperature.

CuSe and SnSe integrate with other phases into CZTSe at ~380 and ~400 C,

respectively, resulting in a kesterite structure which is detected at ~410 C. After peaks h

(CuSe) vanish at ~380 C, peak d (SnSe) disappears at ~400 C. Afterwards,

distinguishable reflections of CZTSe (peaks o) become visible at ~410 °C (Figure 4.14).

The overall reaction path of sample #3-1 is simply described as an arrow diagram in

the Figure 4.16. The rectangles with a gradient green color signify the gradual increase of

its phase from the beginning of rectangles, i.e. the reflection of ZnSe appears from

~330 °C and SnSe from ~350 °C. The rectangles with a gradient red color denote the

rapid decrease of Se reflections at ~210 °C and complete disappearance at ~220 °C.

Because the formation temperature of CTSe is uncertain and because CTSe generally

forms from CuSe and SnSe phases, CTSe is marked by grey color on the arrows of CuSe

Figure 4.16: An arrow diagram for the reaction pathway of a Mo/Cu/Sn/Zn/Se (sample #3-1) stacked layer during heat treatment. The rectangles with gradient color denote the decreasing (faint color) and increasing (vivid color) reflections of each phase.


56

and SnSe. Contrary to other samples, Cu–Se compounds in this sample do not appear

below ~350 C. Only the Cu6Sn5 phase appears, although Cu prefers not to react with Sn

in comparison with Se: Cu–Sn < Cu–Se.

4.2.1.2 Reactions of Mo/Zn/Sn/Se/Cu

In the precursor of sample #3-2 (Mo/Zn/Sn/Se/Cu), Zn (peaks a), Sn (peaks c) and

CuSe (peak h) are obviously observed, as shown in Figure 4.17 at ~30 C. The absence of

any metallic alloys in this precursor is caused by the separation of Cu from Zn and Sn by

the Se layer. This means that peak Z, which indicates a reflection at ~43°, denotes Zn

and/or Cu without any metallic alloys: as mentioned above, peak Z can denotes not only

Cu and Zn but also Cu6Sn5 and Cu–Zn alloys. However, the peak Z in this case seems to

denote only Zn, since its intensity does not change with other reflections of Cu–Se alloys

(peaks h and j) and instead vanishes along with the other Zn reflections (peaks a).

Therefore, peak Z here is also marked by ‘a’ to indicate only the Zn reflection. The

undetectable Cu reflection in Figure 4.17 signifies that most of Cu has combined with Se,

forming CuSe as the Cu is sputtered on the Se layer. Thus, peak h (CuSe) is clearly

observed in the precursor.

Because of the combination of most of the Cu with Se, Se reflections appear at

~140 C higher than that for other cases in general (~110 C). When the Se layer deposits

Figure 4.17: Time–temperature evolution of powder diffractograms for sample #3-2 (Mo/Zn/Sn/Se/Cu) with colour-coded intensities (a.u.) of Bragg reflections. Phases denoted by peaks a-Z are as follows: a: Zn; b: ZnSe; c: Sn; d: SnSe; h: CuSe; j: CuSe2; K: ZnSe (b) and Cu2ZnSnSe4 (o), o: Cu2ZnSnSe4, Z: Zn (a) and possibly Cu.


57

on the Cu layer, as in the case of sample #2-2 (Mo/Sn/Cu/Se) or #2-3 (Mo/Zn/Cu/Se), Se

reflections emerge at ~110 C, as shown in Figure 4.5 or 4.7. This differ from Figure 4.17,

which presents the emergence of Se reflections at ~140 °C. This difference implies that

the crystallisation temperature of Se is influenced by the stacking order of Cu and Se

layers in the precursor. Only when the Cu layer is deposited on the Se layer (Mo/metals/

Se/Cu), the crystallization temperature of Se seems to be delayed. This reason can infer

from the reaction path of sample #2-2. This previous result shows the appearance of Sn

reflections after Cu leaches from Cu6Sn5 (Figure 4.5). Here, the reaction seems to be

similar to the replacement of Sn layer by Se layer. Similar to the revealed metallic Sn

structure in sample #2-2 by Cu diffusion from the Cu–Sn alloy, Se structure in sample #3-

2 is also able to crystallise after Cu leaches from CuSe alloy by outward diffusion of Cu.

Consequently, Se reflections are detectable at ~140 C. Afterward the Se reflections

suddenly weaken at ~200 C and completely disappear at ~210 C because of the use of

Se for the formation of CuSe2.

At ~200 °C, CuSe transforms into CuSe2, and at the same time, the eutectic Sn–Zn

alloy forms, causing the disappearance of Sn reflections. As the temperature rises, the Sn

reflections (peaks c) over the crystallised Se gradually weaken (Figure 4.17), similar to

the results for samples #2-5 and #2-6. However, the reason for this is unveiled in this

study. Afterwards, peaks c (Sn) disappear at ~200 C, indicating the formation of a

eutectic Sn–Zn alloy, in accordance with the results for samples #2-5 and #2-6. Because

only a relatively small amount of Zn is used at the beginning of formation of the Sn–Zn

alloy, the steadily detectable Zn reflections (peak a) after disappearance of Sn reflections

(peak c) is understandable, as revealed in sample #2-5. While the Sn–Zn alloy forms at the

bottom of the film, CuSe2 simultaneously forms from CuSe and crystalline Se on the top

of the film. Consequently, peak h (CuSe) and Se reflections disappear at ~200 and

~210 C, respectively, as soon as peaks j (CuSe2) emerge. This early disappearance of Se

reflections is understandable by the consumption of Se in the formation of the CuSe2

phase at this temperature. Additionally, the result for sample #2-3 reveals that Cu can

react not only with liquid/gas Se but also with crystalline Se structure (see Figure 4.7).

That is, the combination of CuSe with crystalline Se induces the formation of CuSe2 and

disappearance of Se reflections at temperatures lower than its melting point.

At ~290 C, ZnSe is compounded by Se diffusing from CuSe2 and Zn from the molten

eutectic Sn–Zn alloy. As soon as peaks a (Zn) fade, peak b (the beginning of peak K)

emerges in Figure 4.17 at this temperature. Although the Sn layer is in contact with the Se

layer in the precursor (Mo/Zn/Sn/Se/Cu), Sn does not react with Se before Zn reacts. This

formation sequence is same as that for sample #2-6 (Mo/Zn/Sn/Se). Because of the

formation of a eutectic Sn–Zn alloy beneath the Se layer in sample #2-6 and because of

the stronger affinity of Se to Zn than to Sn [22, 23], ZnSe forms earlier than SnSe does


58

after the melting of a Sn–Zn mixture, regardless of the sequence of Sn and Zn layers

under the Se layer. In the same manner, ZnSe in sample #3-3 forms as soon as the Sn–Zn

mixture becomes a liquid phase at ~290 °C. This liquidus temperature indicates that the

eutectic Sn–Zn alloy consists of ~65 at% of Sn and ~35 at% of Zn (Sn-35 at% Zn) in

accordance with the Sn–Zn phase diagram [16] (see section 2.1.1.). As the Zn:Sn ratio in

the initial precursor of sample #3-2 is 1.2:1, the maximum Sn proportion in the eutectic

Sn-Zn alloy is ~45 at% (Sn-54 at% Zn). This composition of the eutectic Sn–Zn alloy (Sn-

54 at% Zn) in terms of the elemental ratio of the initial precursor of sample #3-2 is

completely different from the composition (Sn-35 at% Zn) in terms of the formation

temperature of ZnSe from the in situ analysis. This difference signifies the variation of the

distribution of Sn within the eutectic Sn–Zn alloy, ie, a large amount of Sn on the upper

part of Sn–Zn alloy and a small amount of Sn on the lower part of Sn-Zn alloy. This

declining Sn concentration from top to bottom of the Sn–Zn alloy is understandable,

considering the penetration of Sn into Zn layer during the formation of eutectic alloy and

the sequence of stacked layers in sample #3-2 (Mo/Zn/Sn/Se/Cu). The result for sample

#2-6 (Mo/Zn/Sn/Se) also proves the diffusion of Sn into Zn, forming a eutectic alloy (see

section 4.1.2.6), although ZnSe forms at different temperature. Therefore, the Sn–Zn alloy

melts at ~290 C and Zn reacts with Se at ~290 C because of the lower concentration of

Zn in the eutectic Sn–Zn alloy. The Se for the ZnSe formation derives from CuSe2 by Se

diffusion into the film, as mentioned in section 4.1.3.6. It seems that CuSe2 in sample #3-2

also facillitates the reaction between Se and Zn from the Sn–Zn alloy, similar to the result

for sample #2-3 (Mo/Zn/Cu/Se). Peak b (the beginning of peak K) for another phase such

as CTSe can be interpreted. However, the simultaneous disappearance of peaks a and

appearance of peak K proves that the beginning of peak K denotes ZnSe.

At ~330 C, CuSe and SnSe form at the same time by peritectic decomposition of

CuSe2. When the temperature reaches 332 C, CuSe2 decomposes into CuSe and Se, in

accordance with the Cu–Se phase diagram [18]. The dealloyed Se then reacts with Sn

through equation (8):

CuSe2 + Sn(l) → CuSe + SnSe (~330 C) (8)

Therefore, the strong peaks h (CuSe) and peaks d (SnSe) simultaneously emerge at ~330

C as soon as peaks j (CuSe2) completely vanish. After peaks d (SnSe) fade, peaks o

(CZTSe) become visible in the in situ XRD diffractograms obtained at ~420 °C: the

formation of CZTSe structure at ~420 °C.

The reaction path of sample #3-2 is simply described in Figure 4.18. The rectangle

with gradient green colors indicates that the ZnSe reflection increases gradually from

~300 °C and rapidly from ~350 °C. The other rectangle with gradient red color indicates


59

that the Se reflections decrease rapidly at ~200 °C and disappear completely by ~210 °C.

Because the formation temperature of CTSe is uncertain and because CTSe generally

forms from CuSe and SnSe, CTSe is marked by grey color on the arrows of CuSe and

SnSe.

Figure 4.18: An arrow diagram for the reaction pathway of a Mo/Zn/Sn/Se/Cu (sample #3-2) stacked layer during heat treatment. The rectangles with gradient color denote the decreasing (faint color) and increasing (vivid color) reflections of each phase.

4.2.1.3 Reactions of Mo/Zn/Sn/Cu/Se

Sn, CuSe, CuZn, and traces of Cu5Zn8 and Zn are detected in the initial precursor in

sample #3-3 (Mo/Zn/Sn/Cu/Se). Sn (peaks c) and CuSe (peak h) can be easily observed in

Figure 4.19 at ~30 C, and traces of Cu5Zn8 and Zn are confirmed by ex situ analysis

(GIXD), as described in Figure 4.20. In contrast to Figure 4.15, which is also the result of

ex situ analysis for sample #3-1, only Figure 4.20 evidently presents two weak reflections

of Zn. This indicates that a trace amount of pure Zn element remains in the precursor of

sample #3-3 in contrast to the sample #3-1. The trace amounts of Zn and Cu5Zn8 indicates

that most of Zn combines with Cu, forming β’-CuZn. Although the Sn layer is in contact

with Cu layer (Mo/Zn/Sn/Cu/Se), the Cu–Sn alloy is not detected, similar to the case of

sample #3-1. Non-formation of the Cu–Sn alloy, in contrast to the formation of Cu–Zn

and Cu–Se alloys, is in accordance with the reaction tendency described in section

4.1.3.10: Cu–Sn < Cu–Se < Cu–Zn.

As the temperature rises, CuSe transforms into CuSe2 at ~200 C. Accordingly, the

CuSe2 reflections (peak j) emerge as soon as Se reflections disappear at ~200 C, which is

lower than its melting point. Here the vanishing temperature of Se reflections at ~200 °C

is indubitable because of the simultaneous disappearance of Sn reflections. As observed in

the results for sample #2-5 and #2-6 (see section 4.1.2.5 and 4.1.2.6), Sn reflections may


60

Figure 4.19: Time–temperature evolution of powder diffractograms for sample #3-3 (Mo/Zn/Sn/Cu/Se) with colour-coded intensities (a.u.) of Bragg reflections. Phases denoted by peaks a-Z are as follows: c: Sn, d: SnSe, h: CuSe, j: CuSe2, k: Cu2SnSe3, K: ZnSe, Cu2SnSe3, and Cu2ZnSnSe4 (o), o: Cu2ZnSnSe4, Z: Cu, Zn, Cu5Zn8, CuZn, and α-brass (Cu2Zn).

Figure 4.20: GIXD diffractogram of sample #3-3 (Mo/Zn/Sn/Cu/Se) before heat treatment. This analysis is performed in the atmosphere at room temperature. Trace amounts of Cu5Zn8 and Zn in the precursor are observed without any trace of Zn [72].

vanish at ~200 °C due to the formation of eutectic Sn–Zn alloy. Furthermore, sample #2-2

(Mo/Sn/Cu/Se) and #2-3 (Mo/Zn/Cu/Se) show the Cu diffusion through the Se structure,

causing the earlier disappearance of Se reflections along with the growth of reflections of

the Cu–Se alloy. Thus, Se reflections may disappear before its melting point because of

the Cu diffusion through crystallins Se. The l-CuSe2 reflections (peaks i), which is

observed in Figure 4.5 (sample #2-2) and 4.7 (sample 2-3), are undetectable in Figure 4.19,

because the used amount of Cu for Cu–Se alloy in sample #3-3 is much less than that for


61

two previous samples due to the formation of Cu–Zn alloy in sample #3-3.

Consequently, the eutectic Sn–Zn alloy forms at ~200 C from metallic Sn and Zn

dealloyed from CuZn, while β’-CuZn gradually transforms into Cu2Zn by dezincification.

As mentioned above, the formation of Sn–Zn mixture is comfirmed by disappearance of

Sn reflections (peaks c) at ~200 °C in Figure 4.19. The transformation of β’-CuZn into

Cu2Zn is verified by a gradual shift of peak Z toward lower Bragg angles. The result for

sample #2-4 (Mo/Cu/Zn/Se) reveals that the shifting of peak Z is caused by the

transformation of β’-CuZn into α-brass on the upper side of Cu–Zn alloy by

dezincification. However, the end of the shift of peak Z in Figure 4.19 (sample #3-3) has a

Bragg angle much smaller than that for Figure 4.9 (sample #2-4). The end of peak Z in

Figure 4.19 clearly denotes Cu2Zn, whereas the end of peak Z in Figure 4.9 indicates the

co-existence of Cu2Zn and CuZn (see Figure 4.10). The reason of these different

compositions of Cu–Zn alloys in two samples seems to be the presence or absence of Sn–

Zn alloy inside the film. Zn dealloyed from β’-CuZn in sample #2-4 may not combine

with other elements because sample #2-4 consists of Cu, Zn and Se layers only and

because Cu interrupts the reaction between Zn and Se. Accordingly, Zn moves to the back

electrode, forming Zn-rich CuZn or ε-brass phases on the lower side of Cu–Zn alloy (see

section 4.1.3.4). On the contrary, Zn leaching from CuZn in sample #3-3 may react with

metallic Sn, forming a eutectic Sn–Zn alloy. Because metallic Sn absorbs Zn decomposed

from β’-CuZn, Zn does not recombine with Cu, hence the absence of Zn-rich Cu–Zn alloy

in sample #3-3 in contrast with sample #2-4. For this reason, the end of peak Z in Figure

4.19 can clearly denote only Cu2Zn. It may also recognize that the Cu2Zn forms on the

upper side of Cu–Zn alloy under CuSe2 by outward diffusion of Cu (see section 4.1.3.4).

The temperature of CTSe formation (~300 °C), which will be explained in more detail

about its formation in the succeeding paragraph, also verifies this formation of the eutectic

Sn–Zn alloy in the film. As observed in samples #2-6 and #3-2, the ZnSe forms after the

Sn–Zn alloy becomes a liquid phase. Similar to the ZnSe formation in two previous

samples, here the formation temperature of CTSe in sample #3-3 also corresponds to the

liquidus temperature of Sn–Zn alloy, representing the presence of Sn–Zn mixture in the

film. According to the Sn–Zn phase diagram [16], the eutectic Sn–Zn alloy becomes a

liquid phase at ~300 C when the Sn concentration is near ~62 at% (Sn-38 at% Zn) (see

section 2.1.1.3). Similarly, the maximum Sn concentration of the Sn–Zn mixture in

sample #3-3 reaches this value according to the estimation as follows. To estimate the

proportion of Sn in the eutectic Sn–Zn mixture, it is necessary to adjust the elemental ratio

of the precursor (Cu:Zn:Sn:Se = 1.8:1.2:1:5.3) with respect to the reaction. On the basis of

the phases detected in Figure 4.19 from room temperature to ~200 C, the reaction process

can roughly be described from the preparation of precursor by equation (9).


62

18 Cu + 12 Zn + 10 Sn + 53 Se → 12 CuZn + 6 CuSe + 10 Sn + 47 Se (< ~200 C) (9)

Because Cu prefers to combine with Zn than with Se, in accordance with the reaction

tendency (Cu–Se < Cu–Zn, see section 4.1.3.10), Cu tends to react with Zn, and then the

remaining Cu reacts with Se, forming 12 β’-CuZn and 6 CuSe, proportionally. Afterwards,

CuSe and Se react via equation (6), which is described with the reaction of the Cu–Se

alloys in section 4.1.3.5. Only CuZn and Sn phases become components for the formation

of Sn–Zn mixture, as Zn selectively leaches from CuZn. Because the end of the shift of

peak Z obviously denotes Cu2Zn, the next reaction of β’-CuZn and Sn can roughly be

described by equation (10):

12 CuZn + 10 Sn → 6 Cu2Zn + 6 Zn + 10 Sn (use of Sn–Zn alloy) (10)

At this point, this 6 dealloying Zn and 10 Sn combine to form a Sn–Zn mixture, so that the

maximum concentration of Sn in the Sn–Zn mixture is ~62.5 at% which is near ~62 at%,

as expected. Thus, the liquidus temperature of Sn–Zn alloy is near ~300 C. This estimate

proves the presence of eutectic Sn–Zn mixture in the film and the reason of CTSe

formation at ~300 °C.

At ~300 °C, CTSe forms from CuSe2 and Sn at a slow rate as soon as the Sn–Zn

mixture becomes a liquid phase, as mentioned above. The slowly increasing intensity of

CTSe reflection (peak k) at 300–330 °C in Figure 4.19 is in accordance with the previous

results for sample #2-1 and #2-2 (see section 4.1.3.9). The scarce diminishing of peaks j

(CuSe2) after emergence of peak k at ~300 °C also verifies the CTSe formation at this

stage. In fact, peak K which is marked with small ‘k’ in Figure 4.19 can indicate not only

CTSe but also ZnSe and CZTSe, as mentioned before. However, during measurements of

the three samples (#3-1, #3-2 and #3-3) by in situ XRD, the beginning of peak K for

sample #3-3 emerges at a Bragg angle slightly higher than that for other K peaks of

samples #3-1 and #3-2. These two different Bragg angles are perceptible and are

comparable to other peaks that clearly denote ZnSe and CTSe from samples #2-5

(Mo/Sn/Zn/Se) and #2-1 (Mo/Cu/Sn/Se), respectively. Comparison of these reflections in

Figure 4.21 suggests that the beginning of peak K in sample #3-3 belongs to CTSe,

whereas the beginning of peak K in sample #3-1 belongs to ZnSe. For this reason, the

beginning of peak K is marked ‘k’ to denote the CTSe phase.

One comparable reaction process with this CTSe formation in sample #3-3 is the ZnSe

formation in sample #3-2 (Mo/Zn/Sn/Se/Cu). Although two same phases, CuSe2 and

molten Sn–Zn mixture, are compounded at similar temperature, sample #3-2 forms ZnSe

at ~290 °C whereas sample #3-3 produces CTSe at ~300 °C. According to the reaction

path for sample #3-2, only Sn–Zn mixture was under the CuSe2 before the ZnSe formation.


63

Figure 4.21: Peak K magnified from Figures 4.14 (red) and 4.19 (blue), together with the reflections of ZnSe and Cu2SnSe3 from Figures 4.12 (green) and 4.4 (orange), respectively. These reflections belong to samples #3-1 (Mo/Cu/Sn/Zn/Se; red), #3-3 (Mo/Zn/Sn/Cu/Se; blue), #2-5 (Mo/Sn/Zn/Se; green), and #2-1 (Mo/Cu/Sn/Se; orange).

As Sn-Zn mixture melts at ~290 °C from the upper side, Se diffusing from CuSe2 reacts

with Zn in accordance with the reaction tendencies (Sn–Se < Zn–Se, see section 4.1.3.10).

On the contrary, CuSe2 in sample #3-3 reacts with Sn from the molten Sn–Zn mixture,

forming a CTSe phase. In this case, Se does not react with Zn from the liquid Sn–Zn alloy

although Se diffuses into the film from CuSe2 (see section 4.1.3.5). It signifies that a

certain element in sample #3-3 interrupts the reaction between Zn and Se, and it is Cu

from the Cu–Zn alloy according to the section 4.1.3.7. This previous section shows that

the formation temperature of ZnSe may increase as the Cu concentration in Cu–Zn alloy is

getting higher. In particular, when Cu2Zn forms near the Se, Zn may react with Se at

~360 °C. Sample #3-3 also forms Cu2Zn on the upper side of Cu–Zn alloy under the

CuSe2 because of the metallic Sn, causing the formation of eutectic Sn-Zn alloy. In

addition, Cu is constantly decomposed from the Cu–Zn alloy by dezincification. For these

reasons, Zn from the molten Sn–Zn mixture at ~300 °C prefers to adhere to the Cu–Zn

alloy rather than react with Se; hence, no ZnSe forms in the film at this temperature.

Meanwhile, CuSe2 is able to react with liquid Sn, forming CTSe, similar to the results for

sample #2-1 and #2-2. Therefore, the difference in formation process from the same alloy

phases between samples #3-2 and #3-3 is caused by the presence or absence of Cu–Zn

alloy under the CuSe2.

At ~320 °C, Cu2Zn forms on the upper side of Cu–Zn alloy under CuSe2 because Zn

leaching from Cu–Zn alloy after the CTSe formation remains and moves into the lower

side of the Cu–Zn alloy. Consequently, the shift of peak Z stops at ~320 C and stays at

this Bragg angle, as shown in Figure 4.19. The halted shift of peak Z may be interpreted

as the maintenance of Zn concentration in the Cu–Zn alloy, in accordance with the result

for sample #2-4 (Mo/Cu/Zn/Se). Peak Z in Figure 4.9 (sample #2-4) for this previous

sample is non-shifting at ~230 C and eventually decomposes at ~360 C because Zn


64

could not separate from the Cu–Zn alloy in this sample. This previous result confirms that

stabilisation of the Bragg angle of peak Z signifies a constant Zn concentration in the Cu–

Zn alloy duing dezincification. Because the Sn–Zn mixture melts at ~300 °C in sample

#3-3 via CTSe formation, Zn dealloyed from Cu–Zn alloy cannot be consumed by Sn

anymore. Thus, the Zn migrates to the bottom of the Cu–Zn alloy by outward diffusion of

Cu (see section 4.1.3.4), resulting in a constant proportion of Zn in the Cu–Zn alloy.

At ~330 C, all of CuSe2 decomposes into CuSe becaus the peritectic decomposition

temperature of CuSe2 is 332 C [18]. This reaction path is seen by appearance of peaks h

(CuSe) simultaneously with the disappearance of peaks j (CuSe2) near ~330 °C in Figure

4.19. As revealed in section 4.1.3.9, the rate of CTSe formation varies with the Cu–Se

alloys: the formation rate of CTSe by CuSe and SnSe is faster than that by CuSe2 and

SnSe. According to this previous result, the rate of CTSe formation in sample #3-3 must

increase at this temperature, but the growth rate of peak k in Figure 4.19 does not seem to

be significantly changed in comparison with the strengthening peak k in Figure 4.4

(sample #2-1) and Figure 4.5 (sample #2-2). These two previous Figures showed a rapid

increase of peak k as soon as s-CuSe2 disappears at ~330 °C. However, the peak k in

Figure 4.19 strengthens gradually from ~300 C to ~350 °C until peak Z suddenly

weakens at ~350 C. This seems to be caused by liquid Sn. Because Cu–Se alloys with

SnSe without liquid Sn confirms the difference in formation rate of CTSe in section

4.1.3.9, the reaction of Cu–Se alloys with elemental Sn may be different from the reaction

of Cu–Se alloys with SnSe, as shown in here. It implies that the reaction of CuSe with

SnSe is much easier than with Sn to form CTSe; thus, the formation of CTSe rate does not

change although the CuSe2 decomposes into CuSe.

At ~350 °C, Cu2Zn decomposes and forms ZnSe and Cu2Se phases with diffusing Se,

and simultaneously, the growth rate of CTSe increases as SnSe appears. As shown in

Figure 4.19, peak d (SnSe) is distinctly observed in between two peaks h at around 29–31°,

indicating the formation of SnSe at ~350 °C. According to the reaction tendency described

in section 4.1.3.10, Se may react with Sn after combining with Zn and Cu (Sn–Se < Cu–

Se < Zn–Se). It signifies that ZnSe had already formed before the SnSe formation. The

formation of ZnSe can be confirmed by disappearance of peak Z in accordance with the

result for sample #2-4. This previous result shows the ZnSe formation as soon as Cu–Zn

alloys decompose. Similar to this, peak Z in Figure 4.19 also weakens and disappears at

~350 and ~360 °C, respectively, representing the decomposition of Cu2Zn. Furthermore,

peak K in Figure 4.19 shifts slightly to the left which may be caused by the ZnSe

formation as shown in Figure 4.21: the ZnSe reflection has smaller Bragg angle than the

CTSe reflection has. Therefore, ZnSe formation at this temperature is obvious because of

the disappearance of peak Z (Cu2Zn) and the appearance of peak d (SnSe). While ZnSe

forms by the decomposition of Cu2Zn, the rest of Cu decomposed from Cu2Zn also


65

combines with Se, forming Cu2Se at the same time. Therefore, peak l (Cu2Se) emerges as

soon as peak Z diminishes, but soon disappears thereafter at ~360 C. This sudden

disappearance of Cu2Se indicates that, it combines with SnSe for CTSe formation as soon

as Cu2Se forms, similar to the results for sample #2-1 and #2-2. It also signifies that the

Cu decomposed from Cu2Zn prefers to combine with SnSe and Se rather than form Cu2Se.

Additionally, the rate of CTSe formation becomes higher because of the SnSe formation.

As mentioned in the above paragraph, the rate of CTSe formation may vary depending on

the Sn and SnSe phases and depending on the Cu–Se alloys. Because of the presence of

SnSe and CuSe phases at this temperature (~350 °C), CTSe may grow at a fast rate.

After SnSe integrates into CZTSe at ~400 C, the CZTSe kestertie structure appears.

After peak d (SnSe) disappears completely in Figure 4.19, distinguishable reflections of

CZTSe (peaks o) emerge at ~400 °C, corresponding to formation of the CZTSe kesterite

structure.

All reaction pathways of sample #3-3 are described in Figure 4.22 as well by the

arrow diagram. The rectangle with gradient green colors indicates that the CTSe reflection

increases gradually from ~300 °C and rapidly from ~350 °C. Because Se reflecions

disappear immediately at ~200 °C without any weekening, the red rectengular is not used

in Figure 4.22. The arrow from CuSe pointing at the CTSe through SnSe indicates the

formation of CTSe by CuSe and SnSe. The arrow from Sn–Zn to Cu2Zn expresses the

recombination of Zn with Cu–Zn alloy after the melting of eutectic Sn–Zn alloy. Here

also the peritectic decomposition temperature of CuSe2 is marked by blue botted line at

330 °C.

Figure 4.22: An arrow diagram for the reaction pathway of a Mo/Zn/Sn/Cu/Se (sample #3-3) stacked layer during heat treatment. The rectangles with gradient color denote the decreasing (faint color) and increasing (vivid color) reflections of each phase. The reaction path for sample #3-3 (Mo/Zn/Sn/Cu/Se).


66

4.2.1.4 Detection of residual ZnSe by Raman scattering

As stated in section 3.2.1 (see Figure 3.2), two sandwiched pieces used for the in situ

measurement of one sample are taken from the sample holder after thermal treatment and

analysed by Raman spectroscopy. These two pieces show different amounts of alloy

phases according to Raman scattering from different surface shapes, as described in

Figure 4.23. The optical images on the left upper and bottom sides of Figure 4.23 belong

to each piece after the in situ measurement of sample #3-2. Two coloured Raman

scatterings on the right side of Figure 4.23 corresponds to each coloured area marked by

square and points on the left optical images. As shown in here, the different amount of

ZnSe in each piece is detected although these two pieces were obtained from the same

precursor and then measured at the same time. For this reason, the average of Raman

spectra is calculated according to the pieces and not according to the samples.

Figure 4.23: Plane sectional microscopy images (left) of sample #3-2 (Mo/Zn/Sn/Se/Cu) after heat treatment and correlated Raman scattering results (right) in different zones [72].

To estimate the amount of residual ZnSe, Raman spectra is subjected to peak fitting,

and the relative intensity of major spectra of CZTSe (193–195 cm−1

) is calculated from

the major spectra of ZnSe (246–248 cm−1

) in Table 4.2. A larger number denotes a lower

amount of ZnSe (ICZTSe:IZnSe = #:1) and vice versa. Results for the two pieces from each

sample have designations ‘Upper piece’ (top row) and ‘Lower piece’ (bottom row) in

Table 4.2, but these designations have no relation to the depth profile of the film.

Table 4.2 shows similar results for each piece from samples #3-1 and #3-2. One part

(top row) of these two samples includes levels of the ZnSe phase far greater than that of

the other parts (bottom row). This indicates varying proportions of ZnSe with the depth of

the film. Contrary to this, sample #3-3 appears to have an even distribution over the film,

with parts having similar ratios. However, it has a slightly large amount of the ZnSe phase

in comparison with the other parts (bottom row) of samples #3-1 and 3-2. In general,

samples #3-1 and #3-2 have similarly large amounts of ZnSe with uneven distribution,

whereas sample #3-3 has a lower amount but with even distribution.


67

Table 4.2: Relative intensities of Raman spectra for Cu2Zn SnSe4 at 193–195 cm−1

based on the major spectra of ZnSe at 246–248 cm

−1. The ratios indicate intensities

of CZTSe relative to the intensity of ZnSe (ICZTSe:IZnSe). Data for the one (Upper) and other (Lower) pieces of each of the three samples are shown [72].

Sample #3-1 Sample #3-2 Sample #3-3

Upper piece 3.7:1 4.5:1 8.3:1

Lower piece 10.2:1 10.0:1 8.7:1

4.2.1.5 Discussion

This section 4.2.1 starts with one hypothesis: the residual amount of ZnSe in the film

is related with the sequence of formation of components for CZTSe formation. Regarding

the formation of kesterite structure, two equations are well known as follows.

Cu2SnSe3 + ZnSe = Cu2ZnSnSe4 (11)

Cu2Se + ZnSe + SnSe2 (or SnSe + 1/2 Se2) = Cu2ZnSnSe4 (12)

These equations (11) and (12) show that there are five components of CZTSe: CTSe

for equation (11), Cu2Se and SnSe or SnSe2 for equation (12), and ZnSe for both

equations. Only ZnSe is used in both formation reactions. The reaction paths for samples

#3-1, #3-2, and #3-3 indicate that the formation reactions of CZTSe in these three samples

typically follow equation (11), as evidenced by the undetectable Cu2Se reflection in

Figures 4.14, 4.17 and 4.19. Therefore, we focus on the sequence of ZnSe and CTSe

formations and then compare it to the residual amount of ZnSe in the film. The relative

amounts of remaining ZnSe in each sample are already observed in previous section

Table 4.3: The observed formation temperatures of each selenide in three-metal samples. The unit for the temperature number is the degree Celsius [°C]. Depending on the sequence of elemental stacking layers in the precursor and the amount of Zn in contact with Se, the formation temperatures of ZnSe and Cu2SnSe3 are changed.

Three metallic layers Sample

Selenide

#3-1 #3-2 #3-3

Mo/Cu/Sn/Zn/Se Mo/Zn/Sn/Se/Cu Mo/Zn/Sn/Cu/Se

ZnSe ~330 ~290 ~350

CuSe ~350 ~330 ~330

SnSe ~350 (s) ~330 ~350

CTSe ~350 (p) ~330 (p) ~300

CZTSe ~410 ~420 ~400

(s): relatively low formation rate of SnSe compared with that of CuSe and SnSe for samples #3-1 and #3-2, respectively

(p): probable temperature of alloy formation


68

4.2.1.4. The formation sequence of these two phases can be verified by the formation

temperatures of four phases, including CuSe and SnSe, because CTSe is generally formed

by CuSe and SnSe in these three samples. Thus, the formation temperatures of each

selenides are organised in Table 4.3.

Upon inspection, the sequence of formation between ZnSe and CTSe varies among

samples #3-1, #3-2 and #3-3. ZnSe forms earlier than does CTSe in samples #3-1 and #3-

2, whereas CTSe forms earlier than ZnSe in sample #3-3. CTSe formation is not obvious

in Figure 4.14 (samples #3-1) and Figure 4.17 (sample #3-2), but the formation temper-

ature of CTSe is discernible from the emergence of SnSe reflections. Because CuSe and

SnSe can form CTSe at ~330 C (see section 4.1.3.9) and because SnSe never forms

before CuSe does (see section 4.1.3.10), CTSe formation is possible if the SnSe reflection

is detected together with CuSe reflections at temperatures higher than or equal to ~330 C.

Thus, probable temperatures of CTSe phase formation are marked by ‘p’ in Table 4.3

(unit: degree Celsius). ‘s’ in the column of sample #3-1 indicates the relatively low

formation rate of SnSe in comparison with the growth rate of CuSe in sample #3-1 or with

the growing rate of SnSe in sample #3-2.

According to Table 4.2 for the Raman analysis results, the residual amount of ZnSe in

each sample corresponds to different sequences of ZnSe and CTSe formation. As shown

above, only sample #3-3 forms ZnSe later than it forms CTSe, and only sample #3-3 has a

low amount of residual ZnSe with an even distribution. On the contrary, samples #3-1 and

#3-2, which have relatively large amounts of ZnSe in general, form ZnSe earlier than

CTSe formation. This demonstrates that the amount of remaining ZnSe in the synthesised

kesterite film is variable depending on the sequence of ZnSe and CTSe formations. When

CTSe had formed near Zn and Se in the film and when Zn and Se reach enough energy to

form ZnSe, these two reactive elements (Zn and Se) easily adhere to CTSe, forming

CZTSe. In other words, when Zn and Se can form ZnSe, these two elements prefer to

combine with CTSe to form CZTSe rather than react only each other, thus decreasing the

amount of remaining ZnSe in sample #3-3.

The reason of different formation sequence in sample #3-3 is the co-existence of Cu–

Zn and Sn–Zn alloy under CuSe2 in the film. As described in section 4.2.1.3, although

molten Sn–Zn alloy and CuSe2 were equally in the sample #3-2 and #3-3, sample #3-2

forms ZnSe whereas sample #3-3 makes CTSe as following reaction process:

i) As Zn dealloys from β’-CuZn by dezincification, Zn adheres to metallic Sn, forming

a eutectic Sn–Zn alloy next to the β’-CuZn.

ii) Because Zn concentration in the Cu–Zn alloy steadily decreases by formation of

Sn–Zn mixture, the upper side of Cu–Zn alloy may transforms into Cu2Zn by

outward diffusion of Cu dealloyed from β’-CuZn.

iii) Cu2Zn, precisely the high concentration of Cu on the upper side of Cu–Zn alloy,


69

interrupts the reaction between Se diffusing from CuSe2 and Zn from the Cu–Zn

alloy (see section 4.1.3.7).

iv) When Sn–Zn alloy melts at a liquidus temperature which varies depending on the Sn

concentration [16], Zn from the molten Sn–Zn alloy adheres to Cu–Zn alloy again –

the interruption of Se diffusion by lots amount of Cu on the upper side of Cu–Zn

alloy may be applied to the reaction between Se diffusing from CuSe2 and Zn from

the molten Sn–Zn alloy.

v) Meanwhile, Sn from the molten Sn–Zn alloy reacts with CuSe2, forming CTSe.

Thus, the ZnSe formation is interrupted by Cu2Zn, and this Cu2Zn formation is induced by

the formation of Sn–Zn alloy; the co-existence of Cu–Zn and Sn–Zn alloys under CuSe2

causes the CTSe formation earlier than the ZnSe formation.

Although ZnSe forms later than CTSe in sample #3-3, a low amount of ZnSe is still

detectable in the film. This may be explained by two reasons. One possible reason can be

found in the formation reactions of CTSe and Cu–Zn alloy in section 4.2.1.3. While Cu–

Zn alloy in sample #3-3 undergoes dezincification, dealloyed Cu diffuses into the upper

part of the Cu–Zn alloy and dealloyed Zn forms a eutectic Sn–Zn alloy near Sn. The

sequence of stacked layers in the initial precursor and the island-like Sn layer

(Mo/Zn/Sn/Cu/Se) imply that the Sn–Zn alloy seems to form alongside Cu–Zn alloy under

CuSe2, yielding Mo/Sn–Zn and Cu–Zn/CuSe2. As Sn–Zn alloy becomes a liquid phase at

~300 C, Sn forms CTSe and Zn adheres to the Cu–Zn alloy. Therefore, as CTSe forms

between CuSe2 and the Sn-Zn alloy, the Zn concentration on the lower part of the Cu–Zn

alloy steadily increases because Zn commonly moves to the back electrode by Cu (see

section 4.1.3.4). At that time, some of then ZnSe has enough time to crystallise before

CTSe forms near Zn and Se at the bottom of the Cu–Zn alloy when Se diffusion is faster

than CTSe formation. This explains the residual ZnSe inside the sample #3-3, although in

this study cannot confirm such explanation. Only smooth Se diffusion through Cu is

verified in section 4.1.3.6. As the CuSe2 phase is compounded, Se in the Mo/[metal]/

Cu/Se samples diffuses smooth, resulting in reaction between Se and the metal without

any disturbance. Another possible reason for the presence of residual ZnSe is the Cu-poor

and Zn-rich composition of the initial precursor (Cu:Zn:Sn = 1.8:1.2:1). In comparison

with the stoichiometric ratio of CZTSe (Cu2ZnSnSe4), the initial precursor of three-metal

samples include lots amount of Zn, hence inevitable result of the remaining ZnSe inside

the film. The second reason seems to be discernible reason more than the first reason for

the remaining ZnSe in sample #3-3.

In conclusion, the later formation of ZnSe compared with CTSe formation results in a

low amount of residual ZnSe or its absence in the film. In other words, presence of the

CTSe phase or the Cu2Se and SnSe (or SnSe2) phases near Zn and Se elements may

interrupt the formation of pure ZnSe phase. Furthermore, the co-existence of Cu–Zn and


70

Sn–Zn alloys under CuSe2 (the Mo/Cu–Zn and Sn–Zn/CuSe2 sequence) is the key point to

form CTSe before the ZnSe formation.

4.2.2 Different formation process of two samples with reversed elemental stacking order

In this section, the effect of the sequence direction of stacking layers on the formation

process is investigated by comparison between samples #3-3 and #3-4. As shown above,

sample #3-3 (Mo/Zn/Sn/Cu/Se) forms ZnSe after CTSe formation, causing the smallest

amount of residual ZnS among the three samples, #3-1, #3-2 and #3-3. Therefore, the

formation sequence of ZnSe and CTSe phases is observed again by inverting the sequence

of stacked layers of sample #3-3. Consequently, sample #3-4 has a reversed order of

elemental layers for sample #3-3: the sequence of Mo/Se/Cu/Sn/Zn for sample #3-4.

Because the reaction of sample #3-3 is already described in detail in section 4.2.1.3,

only the reaction path of sample #3-4 (Mo/Se/Cu/Sn/Zn) is explained here. Comparison

between the two samples is described in section 4.2.2.2.

4.2.2.1 Reaction in Mo/Zn/Sn/Cu/Se and Mo/Se/Cu/Sn/Zn with reversed

stacking order

CuSe (peak h), Sn (peaks c) and Zn (peaks a and Z) phases in the initial precursor of

sample #3-4 (Mo/Se/Cu/Sn/Zn) are detected by using an in situ XRD diffractogram, as

described in Figure 4.24. The detectable peak h (CuSe) at room temperature is plausible

because the Cu layer is in contact with the Se layer at the bottom of the precursor and

because Cu is very reactive (see sectoin 4.1.3.2). Here, Cu6Sn5 is not observed in the

precursor in the same manner as that for sample #3-3, although Cu and Sn layers are in

contact with each other. The Cu–Zn alloys (peak Z) are also not detected in Figure 4.24,

differently from sample #3-3. In fact, peak Z can be considered as β’-CuZn and/or Cu

phases, but it denotes only a Zn phase, because the shift of this peak in Figure 4.24 is

similar to that of peak Z in Figure 4.1 (sample #1-1, Mo/Zn/Se). Peak Z in Figure 4.1

which obviously denotes Zn shifts to low Bragg angles, forming a line without any sudden

change of its intensity or Bragg angle. As same as the movement of peak Z in Figure 4.1,

peak Z in Figure 4.24 also does not presents the sudden change of its intensity but moves

gradually to low angles. Additionally, peak Z disappears at the same time as other Zn

reflections (peaks a), hence the denotation of Zn for peak Z. For the same reason, peak Z

does not seem to denote Cu, too, because its intensity does not change while peak h (CuSe)

grows, whereas peak Z in Figure 4.7 (sample #2-3, Mo/Zn/Cu/Se) does. The undetectable


71

Figure 4.24: Time–temperature evolution of powder diffractograms of sample #3-4 (Mo/Se/Cu/Sn/Zn) with colour-coded intensities (a.u.) of the Bragg reflections. Phases denoted by peaks a-Z are as follows: a: Zn, b: ZnSe, c: Sn, d: SnSe, h: CuSe, j: CuSe2, K: ZnSe (b), Cu2SnSe3 and Cu2ZnSnSe4 (o), o: Cu2ZnSnSe4, Z: Zn (a) and possibly Cu.

reflections of Cu6Sn5, CuZn and Cu in the precursor signify that most of the Cu combines

with Se during Cu sputtering on the Se layer.

At ~140 C, the amorphous Se layer crystallises. This crystallisation temperature of Se

is higher than the generally observed temperature, ~110 C, for the same reason for

sample #3-2 (Mo/Zn/Sn/Se/Cu). While the Cu layer is deposited on the Se layer during

sputtering, most of Cu combines with the Se layer. Subsequently, as Cu leaches from the

Cu–Se alloy by outward diffusion of Cu, the remaining Se on the bottom of film can

crystallise as a Se structure. Sample #2-2 (Mo/Sn/Cu/Se) also presents a similar formation

process, as evidenced by the appearance of pure Sn structure from Cu6Sn5 via outward

diffusion of Cu instead of pure Se from the Cu–Se alloy. This implies that the Se

reflections for sample #3-4 do not appear at ~110 C because Se combines with Cu. For

this reason, the high crystallisation temperature of Se can be observed only when the Cu

layer is deposited on the Se layer, e.g. sample #3-2 (Mo/Zn/Sn/Se/Cu).

At ~200 C, CuSe2 is formed by CuSe and crystalline Se on the lower part of the film.

Consequently, Se reflections suddenly weaken at ~200 C and completely disappear at

~210 C, which is lower than its melting point, because of the requirement of a large

amount of Se for CuSe2 formation. This sudden weakening of Se reflections is observed

near ~200 °C with the change of the colour of light from green to blue in Figure 4.24

because the colour indicates the intensity of reflection. At the same time, peak h (CuSe)

suddenly weakens as soon as peaks j (CuSe2) appear at ~200 C. It signifies that much Se

and CuSe are required to form the CuSe2 phase. Therefore, Se reflections disappear below


72

the melting point (221 °C) as same as the result for sample #3-2.

Meanwhile, Sn and Zn combine with each other at ~200 C to form a eutectic Sn–Zn

alloy on the upper part of film. This formation of Sn–Zn alloy may be confirmed by the

disappearance of peaks c (Sn) at ~200 C, in accordance with the results for samples #2-5

(Mo/Sn/Zn/Se) and #2-6 (Mo/Zn/Sn/Se). In particular, the change in intensity of peaks a

(Zn) in Figure 4.24 is similar to that in Figure 4.12 for sample #2-5. One Zn reflection at

~36° for the (002) plane strengthens as soon as Se reflections disappear, whereas the other

Zn reflection at ~43° for a (101) plane weakens while both two reflections shift to lower

Bragg angles during heat treatment. This implies that the Zn structure is very oriented

toward the [002] direction as the crystalline Se disappears and as the eutectic Sn–Zn alloy

gradually forms through the Zn and Sn layers, that is, during Sn diffusion through the Zn

layer. Because of the absorption of all Zn into Sn for the formation of Sn–Zn mixture, Zn

reflections (peaks a and Z) disappear completely at ~260 °C.

At ~270 °C, ZnSe gradually forms, indicating the melting of eutectic Sn–Zn alloy. The

emergence of peak b in Figure 4.24 represents this reaction path at this temperature. In

fact, the peak b which is the beginning of peak K can denote not only ZnSe but also CTSe

and CZTSe, as mentioned before. However, undetectable Cu–Zn alloy near this temper-

ature assures the denotation of ZnSe for the beginning of peak K, according to the result

for sample #3-3. Unless Sn–Zn alloy coexists with Cu–Zn alloy under (or near) CuSe2,

CTSe cannot form before the ZnSe formation, as described in section 4.2.1.5.

Additionally, Se prefers to react with Zn than with Sn (Sn–Se < Zn–Se; see section 4.1.3.

10). For these reasons, sample #3-2 also forms ZnSe at ~290 °C from the molten Sn–Zn

alloy and Se diffusing from CuSe2. Similar to the result for sample #3-2, sample #3-4 also

does not form any of Cu–Zn alloys during the measurement. Therefore the beginning of

peak K in Figure 4.24 is obviously produced by ZnSe (peak b). However, the formation

temperature of ZnSe in sample #3-4 is much lower than previous results described in

Table 4.1. Even sample #1-1 (Mo/Zn/Se) which has only Zn and Se layers in the precursor

forms ZnSe at temperature (~290 °C) higher than that for sample #3-4 (~270 °C). It

signifies that Zn is activated by something, and the answer seems to be the melting of

eutectic Sn–Zn alloy. Because the Sn–Zn alloy becomes liquid phase at ~270 °C, the Zn

can react with Se diffusing from CuSe2, similar to the result for sample #3-2. Sample #3-2

also presents the ZnSe formation after Sn–Zn melts at ~290 °C. Only the difference in

results between sample #3-2 and #3-4 is the liquidus temperature of Sn–Zn alloy because

of the different concentration of Sn in its alloy near Se. According to the measured

liquidus temperature of Sn–Zn alloy in sample #3-4, the Sn–Zn mixture, which has

formed on the CuSe2 (Mo/CuSe2/Sn–Zn), consists of ~71 at% of Sn (Sn-29 at% Zn) [16].

In the same manner, the composition of Sn–Zn mixture under CuSe2 (Mo/Sn–Zn/CuSe2)

in sample #3-2 is Sn-35 at% Zn. Considering the stacking sequence of Sn and Zn layers in


73

two samples, the Mo/Se/Cu/Sn/Zn sequence (sample #3-4) converts into Mo/CuSe2/Sn-29

at% Zn/Sn–Zn at ~270 °C, whereas the Mo/Zn/Sn/Se/Cu sequence (sample #3-2) becomes

Mo/Sn–Zn/Sn-35 at% Zn/CuSe2 at ~290 °C. It suggests that Sn seems to prefer to diffuse

inward rather than outward because the Sn concentration in Sn–Zn alloy near CuSe2 for

sample #3-4 is higher than that for sample #3-2.

At ~300 C, SnSe forms from Se diffusing from CuSe2 and liquid Sn from the molten

Sn–Zn alloy. Consequently, CTSe formation is possible by formation of SnSe and the

presence of CuSe2 at this temperature. As shown in Figure 4.24, peak d (SnSe) appears at

~300 C while peaks j (CuSe2) appear, in contrast to Figures 4.14 (sample #3-1), 4.17

(sample #3-2) and 4.19 (sample #3-3), which present the appearance of peak d after the

disappearance of peaks j. The formation temperature of SnSe in sample #3-4 is much

lower than that in other samples because Sn–Zn alloy melts at low temperature, causing

the ZnSe formation. As soon as SnSe forms, CTSe can form at a low rate, in accordance

with equation (2) in section 4.1.2.1. Although Figure 4.24 cannot confirm CTSe formation

because of the overlapping reflections of ZnSe and CTSe at ~27° (peak K), the previous

result for sample #2-1 suggests CTSe formation from CuSe2 and SnSe: 2 CuSe2 + SnSe →

Cu2SnSe3 + Se. For this reason, peaks j (CuSe2) gradually weaken when the intensity of

peak K (ZnSe and CTSe) increases along with the appearance of peak d (SnSe) at ~300 C,

as shown in Figure 4.24. The simultaneous change of their intensities evidences CTSe

formation at this temperature.

The rate of CTSe formation increases at ~330 °C due to the decomposition of CuSe2

into CuSe. Subsequently, SnSe and CuSe disappear at ~340 and ~360 °C, respectively.

According to section 4.1.3.9, the rate of CTSe formation depends on the Cu–Se alloys. It

is revealed that CTSe formation from CuSe and SnSe is faster than that from CuSe2 and

SnSe. Additionally, it is well known that the peritectic decomposition temperature of

CuSe2 is 332 C [18]. Decomposition of CuSe2 into CuSe can be observed from the

disappearance of peaks j (CuSe2) along with the appearance of peaks h (CuSe) in Figure

4.24 at ~330 C. Simultaneously, the intensity of peak K (CTSe and ZnSe) at ~27° rapidly

strengthens. In the meantime, peaks d (SnSe) and h (CuSe) vanish at ~340 and ~360 C,

respectively. This gradual weakening of peak h and disappearance of peak d clearly

indicate the formation of CTSe through equation (3) in section 4.1.2.1: 2 CuSe + SnSe →

Cu2SnSe3.

At ~400 C, CZTSe is clearly observable from the faint reflections of CZTSe near ~35°

and ~36° (peaks o), which indicate the kesterite structure. Although peaks d (SnSe) and h

(CuSe) disappear at ~340 and ~360 C, respectively, peaks o (CZTSe) do not appear

below ~400 C, distinct from other samples showing the appearance of peaks o soon after

the disappearance of peak d. According to the literature [73], the CZTSe phase transforms

from cubic to kesterite structure at ~460 C. This result implies the need of the kesterite


74

structure for a certain thermal energy. In the same manner, the absence of peaks o below

~400 C in Figure 4.24 seems to imply that the kesterite structure of the CZTSe phase

requires a certain amount of thermal energy.


with gradient red colors indicates that the Se reflection decreases rapidly at ~200 °C and

disappear completely at ~210 °C. Because the temperature of CTSe formation is uncertain

and because CTSe may form from CuSe2 and SnSe, CTSe is marked by grey color and is

inserted in between an arrow of CuSe2 and a letter of SnSe.

Figure 4.25: An arrow diagram for the reaction pathway of a Mo/Se/Cu/Sn/Zn (sample #3-4) stacked layer during heat treatment. The rectangles with gradient color denote the decreasing (faint color) and increasing (vivid color) reflections of each phase.

4.2.2.2 Discussion

The reaction path for sample #3-4 (Mo/Se/Cu/Sn/Zn) is different from the result for

sample #3-3 (Mo/Zn/Sn/Cu/Se) although the sequence of stacking layers for sample #3-4

is inversely same with sample #3-3. In particular, sample #3-4 forms ZnSe before the

CTSe formation whereas sample #3-3 forms ZnSe after the CTSe formation. As it

mentioned in section 4.2.1, the different formation sequence of these two phases signifies

a different amount of residual ZnSe in the film. Thus, the change to an inverse sequence

of stacking layers in the precursor may lead to the different amount of residual ZnSe in the

film. To understand the difference in formation sequence of CTSe and ZnSe in spite of the

same stacking order in two initial precursors, it is necessary to observe the resulting alloys

before CTSe and ZnSe form in each sample.

Sample #3-3 forms CTSe earlier than it forms ZnSe. Before CTSe forms, eutectic Sn–

Zn alloy, Cu–Zn alloy and CuSe2 phases are present. As soon as the eutectic Sn–Zn alloy


75

melts at ~300 °C, Sn reacts with CuSe2 while Zn adheres to the Cu–Zn alloy. The reason

of non-combination of Zn and Se at this temperature is the formation of Cu2Zn on the

upper side of Cu–Zn alloy under CuSe2. The formation of ZnSe in this sample occurs by

decomposition of Cu–Zn alloy at ~350 °C. Sample #3-4 forms ZnSe earlier than it forms

CTSe. Before ZnSe forms, CuSe2, eutectic Sn–Zn alloy, metallic Zn and CuSe2 phases are

present. Only one phase of sample #3-3, Cu–Zn alloy, is replaced by metallic Zn in

sample #3-4 because most of Cu is already combined with Se in the initial precursor of

sample #3-4. Accordingly, Se may diffuse from bottom into the film via CuSe2 (see

section 4.1.3.5) because of the absence of Cu2Zn in the film. Therefore, Se may react with

Zn when the Sn–Zn alloy becomes a liquid phase at ~270 °C. The remaining Sn from the

ZnSe formation also reacts with Se at ~300 °C, forming SnSe. Subsequently, CTSe may

form from the SnSe and CuSe2 at ~300 °C.

In any case, CTSe is compounded from the liquid phase of eutectic Sn–Zn alloy. On

the contrary, ZnSe is formed either by melting of Sn–Zn alloy or by decomposition of Cu–

Zn alloy under CuSe2. Depending on the presence or absence of Cu–Zn alloy in the

precursor, the formation sequence of ZnSe and CTSe changes. Therefore the formation of

Cu–Zn alloy in the initial precursor delays ZnSe formation by trapping Zn in the Cu–Zn

alloy, according to observation of these two samples.

In fact, the formation sequence of alloys in sample #3-4 is similar to that of sample

#3-2 (Mo/Zn/Sn/Se/Cu) although the depositing sequence of elemental layers of sample

#3-2 is not the same as that for sample #3-4. Only the formation temperatures of ZnSe,

SnSe and CZTSe phases are different between two samples. Because any of Cu–metal

alloys may not form in the initial precursor in both cases, the compounds in the film are

divided into Sn–Zn and Cu–Se alloys upon heating of these two samples, causing the

same sequence of alloy formations. It signifies that the compounds in the initial precursor

determine the formation sequence of alloys. Therefore the conclusion of this section

becomes similar to one of result described in section 4.2.1.5: the importance of co-

existence of Sn–Zn and Cu–Zn alloys in the film. The comparison between three samples

(#3-2, #3-3 and #3-4) suggests that the consideration of alloys in the precursor is

necessary as much as the stacking order of elemental layers.

4.2.3 The effect of two Cu layers on the reaction

The previous result in section 4.2.1 reveals that ZnSe must be compounded later than

the CTSe formation to reduce or remove residual ZnSe in the film. Furthermore, it is

revealed that the formation of Cu–Zn alloy in the precursor makes delay the ZnSe

formation, as discussed in sections 4.2.1.5 and 4.2.2. On the basis of the previous results,

the Cu layer is divided into two layers and deposited at separate positions, one for the


76

formation of Cu–Zn alloy and the other for the formation of CTSe during annealing. As

shown in the results for sample #2-1 (Mo/Cu/Sn/Se) and #2-2 (Mo/Sn/Cu/Se), the

formation of CTSe is not influenced by stacking order of Cu and Sn layers. Because ZnSe

may form from the molten Sn–Zn alloy at low temperature, Zn and Sn layers are separated

by one of Cu layer. Another Cu layer deposits on the Se layer for sample #3-5 and beneath

the Zn layer for sample #3-6 to minimise the sequence variation of stacked layers in

between two samples. Therefore sample #3-5 has the Mo/Zn/Cu/Sn/Se/Cu sequence, and

sample #3-6 has the Mo/Cu/Zn/Cu/Sn/Se sequence. Consequently, the separate formation

of CTSe and Cu–Zn alloy phases in two samples is expected because CTSe may form at

~290 °C regardless of stacking order of Cu, Sn and Se layers, in accordance with the

results for sample #2-1 and #2-2.

4.2.3.1 Reactions of Mo/Zn/Cu/Sn/Se/Cu

The reaction path of sample #3-5 (Mo/Zn/Cu/Sn/Se/Cu) during the overall measure-

ment is totally different from that of other previous samples, as shown in Figure 4.26.

Even the kesterite structure does not form at the end of its measurement. It seems that the

formation of one new phase in the precursor induces this different result.

The precursor consists of β’-CuZn (peak Z), Cu6Sn5 (peaks f and Z), CuSe (peaks h),

Cu3Se2 (peaks p) and probable Cu5Zn8 (peak Z) phases, as shown in Figure 4.26 at ~30 C.

Because the distinguishable weak reflections of Cu5Zn8 at ~35° and ~38° overlap with

those of Cu3Se2 (peaks p), the presence of Cu5Zn8 in the precursor is uncertain. On the

basis of the alloy phases in the initial precursor of sample #2-3 (Mo/Zn/Cu/Se), which has

the same stacking order as the lower part of sample #3-5 (Mo/Zn/Cu/Sn/Se/Cu), most of

the Cu in sample #3-5 probably combines with Zn, forming the CuZn phase mainly while

Cu is deposited on the Zn layer (see Figure 4.8). Detectable reflections of Cu6Sn5 and

CuSe are also plausible on the basis of the precursors of sample #2-2 (Mo/Sn/Cu/Se). This

previous result shows the formation of Cu6Sn5 and CuSe phases in its precursor because

Cu layer was in contact with Sn and Se layers. In contrast to these phases, Cu3Se2 is a new

compound which has not been observed in this study. According to the sample #3-2

(Mo/Zn/Sn/Se/Cu), the stacking order of Sn/Se/Cu on the upper part of sample #3-5

(Mo/Zn/Cu/Sn/Se/Cu) may produce only CuSe without Cu3Se2 in the initial precursor of

sample #3-5, differently from the observation in Figure 4.26. It signifies that another Cu

layer beneath the Sn layer (Mo/Zn/Cu/Sn/Se/Cu) somehow influences on the formation of

Cu3Se2 in the precursor of sample #3-5. Considering the weakest tendency of Cu to react

with Sn among the other components (Cu–Sn < Cu–Se < Cu–Zn) and the discontinuous

deposition of Sn layer in general [70, 71], it seems that a part of Cu beneath the Sn layer

in sample #3-5 (Mo/Zn/Cu/Sn/Se/Cu) diffuses into the Se layer through the Cu–Sn alloy

while Se and Cu layers are sequentially deposited on the Sn layer. This Cu diffusion from


77

Figure 4.26: Time–temperature evolution of powder diffractograms of sample #3-5 (Mo/Zn/Cu/Sn/Se/Cu) with colour-coded intensities (a.u.) of the Bragg reflections. Phases denoted by peaks a-Z are as follows: b: ZnSe, f: Cu6Sn5, g: Cu3Sn, h: CuSe, l: Cu2–xSe and Cu2Se, n: ε-brass, p: Cu3Se2, q: Cu41Sn11, K: ZnSe (b), Z: Cu, CuZn, Cu6Sn5 (f), CurichZn and Cu3Sn (g).

bottom to top of Se layer seems to be affected by the deposition of not only Se layer but

also Cu layer on the Se layer (Mo/Zn/Cu/Sn/Se/Cu) because this CuSe phase is not

observed in the precursor of sample #2-1 (Mo/Cu/Sn/Se). Additionally, when this Cu

layer (Mo/Zn/Cu/Sn/Se/Cu) is sputtered on the Se layer, most of the Cu combines with Se,

similar to the initial precursor of sample #3-4 (Mo/Se/Cu/Sn/Zn): sample #3-4 shows the

combining of most of the Cu with Se, causing the non formation of any of Cu–metal

alloys in its precursor. Thus, CuSe seems to have formed on the lower part of Se layer by

Cu beneath the Sn layer (Mo/Zn/Cu/Sn/Se/Cu), and another Cu also diffuses from top to

bottom of Se layer during the deposition of the Cu layer (Mo/Zn/Cu/Sn/Se/Cu) on the Se

layer. For this reason, the Cu concentration on the lower part of the Se layer increases, as

much as CuSe transforms into Cu3Se2, in accordance with the Cu-Se phase diagram [18].

Thus, Cu3Se2 (peaks p) forms on the lower part of the Se layer simultaneously with the

formation of CuSe phase (peaks h) on the upper part of the Se layer: Mo/CuZn/Cu6Sn5/

Cu3Se2/CuSe.

At ~110 C, the new compound Cu3Se2 decomposes into the Cu2–xSe phase

(Mo/CuZn/Cu6Sn5/Cu2–xSe/CuSe). At the same time, CuSe grows along the z-axis

direction becasue of Se dealloyed from Cu3Se2. The transformation of Cu3Se2 into Cu2–

xSe is clearly observed from the disappearance of peaks p (Cu3Se2) and the appearance of

peaks l (Cu2–xSe) at ~110 C in Figure 4.26. This temperature is well in accordance with

the decomposition temperature of Cu3Se2 at ~113 °C, as written in Cu–Se phase diagram

[18]. Simultaneous with this transformation, one of the peaks h (CuSe) at ~28° for a (112)


78

plane slightly weaken, while peak h at ~31° for a (006) plane increases. That is the CuSe

structure grows along the [006] direction. This implies that when Cu2–xSe forms via

decomposition of Cu3Se2 on the lower part of the Se layer, which now converts into the

Cu–Se alloy layer, the Se dealloyed from Cu3Se2 somehow affects the CuSe structure on

the upper part of the Se layer to orient along the [006] direction.

As the temperature rises, Cu3Sn (peaks g, Z) forms at ~180 C, and then the amount of

Cu2–xSe alloy (peaks l) gradually grows from ~220 C by Cu diffusion from Cu6Sn5.

Because Cu steadily dealloys from Cu6Sn5 (peaks f, Z) and diffuses into the Cu–Se alloys

by outward diffusion of Cu, peak f (Cu6Sn5) weakens as peaks g (Cu3Sn) emerge along

with the strengthening peak Z (Cu3Sn, β’-CuZn, Cu6Sn5) at ~180 C, indicating the

growth of Cu3Sn. Afterwards, the intensities of peaks l (Cu2–xSe) start to increase at

around 220 °C and steadily strengthen, while the intensities of peaks h (CuSe) decrease

with the strength-ening of peaks l. In this stage, the emerging temperature of Cu3Sn

reflections (peaks g) corresponds to the result of sample #2-1 (Mo/Cu/Sn/Se) at ~180 C.

This previous result shows that Cu3Sn soon becomes depleted and disappears as the

diffusing Cu combines with CuSe and forms CuSe2 at ~220 C (see Figure 4.4). Contrary

to the previous result for sample #2-1, the Cu3Sn phase remains in sample #3-5 until

~410 C. However, the Cu2–xSe phase grows at ~220 C. Because the Cu concentration in

the Cu–Se alloys steadily increases by outward diffusion of Cu from Cu6Sn5 via Cu3Sn,

CuSe can continuously transform into Cu2–xSe (peaks l) after Cu3Sn formation. That

means the gradual conversion of Cu6Sn5 into Cu3Sn and Cu2–xSe, causing a liquid Sn.

Therefore, the inversely changing intensities of peaks l (Cu2–xSe) and peaks h (CuSe) are

observed. The formation of liquid Sn from the decomposition of Cu6Sn5 is plausible in

consideration of the elemental ratio of this sample and the compounding alloys in this

stage. These transformations of alloys suggest that the sequence of alloy layers in the film

can be described as follows: Mo/CuZn/Sn + Cu6Sn5/Cu3Sn/Cu2–xSe/CuSe.

At ~250 C, ε-brass (Cu0.7Zn2) emerges on the lower part of the β’-CuZn phase by

dezincification, and Cu6Sn5 completely decomposes at ~260 C while the amount of Cu2–

xSe gradually increases. As temperature rises, Zn selectively leaches from β’-CuZn and

forms ε-brass on the lower part of Cu–Zn alloy, similar to the result for sample #2-4a.

Consequently peak n (Cu0.7Zn2) slowly emerges with a faint intensity at ~250 °C as peak

Z (β’-CuZn) weakens and shifts to low angles. This peak n can obviously be seen when

the in situ XRD diffractograms are magnified. As shown in the result for sample #2-4a

(see Figure 4.11), the shift of peak n to low Bragg angles indicates an increase in Zn

concentration in the ε-brass phase. On the contrary, the shift of peak Z to low Bragg

angles represents the increase in Cu concentration of β’-CuZn, in accordance with the

result for sample #2-4 (see Figure 4.10). It signifies that Zn moves from upper to lower

part of β’-CuZn alloy by dezincification (see section 4.1.3.4), causing the formations of ε-


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brass and CurichZn alloy on the upper and lower parts of β’-CuZn phase, respectively.

Although Cu also steadily dealloys from β’-CuZn, it does not combine with liquid Sn

decomposed from Cu6Sn5 because of the reaction tendency (Cu–Sn < Cu–Zn; see section

4.1.3.10), thus the presence of Cu3Sn and a liquid Sn under Cu2–xSe (peaks l): Mo/ε-brass/

β’-CuZn/CurichZn/Sn(l)/Cu3Sn/Cu2–xSe/CuSe. The complete decomposition of Cu6Sn5 is

be observed by disappearance of its reflections (peaks f and Z) at ~260 °C in Figure 4.26.

Because the melting point of Sn is 221 °C, the remaining Sn from the decomposition of

Cu6Sn5 would become a liquid phase at this temperature. The formation of a eutectic Sn–

Zn alloy by liquid Sn and dealloyed Zn from Cu–Zn alloy seems to be possible at this

stage. However, the formation is not considered in here because Zn decomposed form β’-

CuZn is generally induced to move into the lower part of Cu–Zn alloy (see section

4.1.3.4).

Afterwards, Cu2–xSe converts into Cu2Se at 290–340 °C because of the increase of its

amount by decomposition of CuSe at ~290 °C, and subsequently the amount of Cu3Sn

slightly increases. Accordingly, peaks l (Cu2–xSe) dramatically shift to low Bragg angles

at 290–340 °C, weakning their intensity, as soon as peaks h (CuSe) disappear at ~290 °C.

When the shift of peaks l halts, peaks g (Cu3Sn) become subtly stronger at ~340 C. This

intends that an acceptable amount of Cu for the Cu-rich Cu–Se alloy is determined. While

Cu diffuses from the Cu–Sn alloys to the Cu–Se alloy, Cu–Se alloy transforms from CuSe

to Cu2Se through Cu2–xSe. As soon as Cu2Se with a certain structural size forms at

~340 C, the amount of Cu reaches the capacity of the Cu-rich Cu–Se alloy (Cu2Se). For

this reason, Cu cannot diffuse into the Cu–Se alloy anymore but induces the small

increase in the amount of Cu3Sn. Additionally, here the Cu causing the slight strength-

ening of Cu3Sn reflections at ~310 °C seems to come from the decomposition of CurichZn

by ZnSe formation.

While Cu2–xSe transforms into Cu2Se, a certain amount of ZnSe forms at ~310 °C by

the decomposition of CurichZn, and the remaining Cu adheres to Cu3Sn. Here, the amount

of ZnSe does not increase as it forms because Se diffusion through film is interrupted by

Cu2Se. Transitions of the reflections in Figure 4.26 suggest this reaction path near this

temperature by peaks K, n and Z. In fact, this peaks K can denote not only ZnSe but also

CTSe and CZTSe, but the result for sample #2-4a verifies the denotation of peak K only

for ZnSe. According to this previous result (see Figure 4.11), ZnSe formation starts when

peak n reaches the lowest Bragg angle and when peak Z disappears. After the appearance

of peak b in Figure 4.11, peak n maintains its lowest Bragg angle for a while and then

shifts to high angles again. Similar to the previous result, peak n in Figure 4.26, which has

shifted to low angles from ~250 °C, also maintains the lowest Bragg angle at 310–340 °C

after the appearance of peak K at ~310 °C. Afterwards, the peak n gradually reverts to

high angles at ~340 °C. The same shift of peak n in both results intends that peak K


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belongs to ZnSe so that the small ‘b’ is written next to the capital K in Figure 4.26 to

indicate only ZnSe without CTSe. One difference between Figures 4.11 and 4.26 is the

non-vanishing peak Z in Figure 4.26 whereas the peak Z in Figure 4.11 disappears after

the emergence of peak b. The reason of it may infer from the non-increasing intensity of

peak b (ZnSe) and the slight increase in intensity of peaks g (Cu3Sn) in Figure 4.26.

Contrary to the strengthening peak b in Figure 4.11, the peak b in Figure 4.26 maintains

its intensity as it appears. This means that the amount of ZnSe does not increase although

the Cu–Zn alloy has enough Zn for ZnSe formation. It also intends that Se does not

diffuse from the Cu–Se alloy through the Cu–Zn alloy. Because of the lack of Se for the

ZnSe formation, the decomposition of Cu–Zn alloy stops and maintains Cu concentration

in the Cu–Zn alloy, causing the non shift and disappearance of peak Z in Figure 4.26. It

signifies that Se may diffuse only when Cu2–xSe decomposes into Cu2Se. In contrast to

facile diffusion of Se through the CuSe2 phase (see section 4.1.3.6), this reaction path

represents that Cu2Se interrupts Se diffusion through the film. Additionally, the presence

of Cu3Sn also influence on the non-vanishing peak Z in Figure 4.26 because Cu3Sn also

produces peak Z. When Figure 4.26 near these temperatures is magnified, the slight

increase in intensity of peaks g (Cu3Sn) at ~310 °C is observed as soon as peak b appears.

It signifies that a certain amount of remaining Cu from the decomposition of CurichZn

alloy by ZnSe formation adheres to Cu3Sn, increasing the amount of Cu3Sn because Cu

may not diffuse into Cu–Se alloy anymore. Therefore, only a certain amount of CurichZn

alloy may decompose and form small amounts of ZnSe and Cu3Sn at this temperature

because Se diffusion is interrupted by Cu2Se.

As all of CurichZn decomposes at ~410 C, the amount of ε-brass increases and Cu3Sn

transforms into Cu41Sn11. It intends that the decomposition of CurichZn induces the

formation of ε-brass and Cu41Sn11. Upon dezincification of the Cu–Zn alloy, the dealloyed

Cu and Zn steadily and respectively move to the upper and lower parts of the Cu–Zn

phases. Therefore the amount of ε-brass increases, and the decomposed Cu adheres to

Cu3Sn, inducing the Cu41Sn11 formation. Accordingly the peak Z (CurichZn) weakens and

disappears at ~410 °C as peak n (ε-brass) grows and reaches highest intensity. In

particular, the strengthening of peak n at 380–410 C is inversely proportional to the

weakening of peak Z. Additionally, peaks g (Cu3Sn) also disappear at ~410 °C along with

the peak Z as faint peaks q (Cu41Sn11) appear at ~34° and ~37°. The main reflection of

Cu41Sn11 at ~42.6° seems to be concealed by peak n because of the small amount of

Cu41Sn11. the decomposition of CurichZn together with the formation of Cu41Sn11 signifies

that the Cu from CurichZn prefers to combine with Sn than Zn at this temperature, in

contrast to the reaction tendency (Cu–Sn < Cu–Zn; section 4.1.3.10). It seems that the

affinity of Cu to Sn becomes stronger than the affinity of Cu to Zn at this temperature

because of the high concentration of Cu in Cu–Zn alloy near liquid Sn. Similarly the


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reaction tendency of Zn to Se and Cu also changes depending on the Cu concentration in

Cu–Zn alloy near Se (see section 4.1.3.10). Without this description, it is difficult to

understand the reaction path at this temperature. Another notable observation at this

temperature is the undetectable SnSe reflection, in contrast to the other general results,

such as those for sample #2-1 (Mo/Cu/Sn/Se). According to the previous result for sample

#2-1, SnSe appears as soon as Cu3Sn disappears (see Figure 4.4). However, any SnSe

reflection is unobservable in Figure 4.26, although Cu3Sn decomposes at this stage,

similar to absence of strengthening of the ZnSe reflection. Here the absence of SnSe peaks

also verifies that Se does not diffuse through the film after Cu2Se forms. As Cu2Se

stabilises, interdiffusion between Se of the Cu–Se alloy and Cu of metallic alloys (Cu–Sn

or Cu–Zn alloys) cease. As Se does not diffuse and thus does not form binary selenides, it

may result in the immiscibility of Cu2Se because Cu2Se can combine only with SnSe (or

SnSe2) and ZnSe phases according to the equation (12) described in section 4.2.1.5. This

result is worthy of note for understanding the remaining Cu2Se in the kesterite film.

Finally, Cu41Sn11 seems to disappear when the sample bursts by evaporation of

gaseous Se and/or SnSe. Therefore, Cu2Se, ZnSe and ε-brass remain in the film after

measurement. The reaction path of sample #3-5 is simply described in Figure 4.27. The

arrow with dotted line from Cu6Sn5 to Cu2–xSe denotes the indirect diffusion of Cu

through Cu3Sn. The faint letters and arrows for Cu41Sn11 and Sn phases mark the

unclearly observed reaction paths in Figure 4.27.

Figure 4.27: An arrow diagram for the reaction pathway of a Mo/Zn/Cu/Sn/Se/Cu (sample #3-5) stacked layer during heat treatment. The faint arrow and letter for Cu41Sn11 at 550 °C denote the undetectable reflections of this phase after the rupture of sample.


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4.2.3.2 Reactions of Mo/Cu/Zn/Cu/Sn/Se

When one Cu layer of sample #3-5 (Mo/Zn/Cu/Sn/Se/Cu) is prepared not on the Se

layer but between Mo and Zn layers, such as a Mo/Cu/Zn/Cu/Sn/Se stacked layer (sample

#3-6), the reaction path for sample #3-6 becomes similar that for other samples, as

described in Figure 4.23. This is in contrast to the reaction path of sample #3-5. In

particular, the formation of alloys at room temperature to ~330 C corresponds to the

results for samples #2-1 (Mo/Cu/Sn/Se) and #2-4a (Mo/Cu/Zn/Se; [Cu]/[Zn] = 1.3).

The precursor of sample #3-6 (Mo/Cu/Zn/Cu/Sn/Se) mainly consists of Cu6Sn5 (peaks

f and Z) and β’-CuZn (peak Z) phases. Thus, only peaks f and Z are detected at ~30 C in

Figure 4.28. Although peak Z can denote Cu6Sn5, β’-CuZn, Cu, Zn and Cu5Zn8, the

distinguishable weak reflections of Zn and Cu5Zn8 are not apparent in Figure 4.28. The

detection of only the β’-CuZn phase without Zn and Cu5Zn8 phases in Figure 4.28 is in

accordance with Figure 4.9 (sample #2-4). The detection of Cu6Sn5 is also similar to the

result for sample #2-1 in Figure 4.4. Metallic Cu, which can also be denoted by peak Z,

does not seem to be present in the precursor. Because Cu is very reactive, as described in

section 4.1.3.2, and because it is divided into two layers in this sample, combination of

most of the Cu with other metallic elements seems to be more plausible than is the

presence of metallic Cu in the precursor. Therefore, the precursor on the lower part and on

the upper part of film is mainly composed of β’-CuZn and Cu6Sn5 phases, respectively,

together with a Se layer on top of the film (Mo/CuZn/Cu6Sn5/Se layer).

Upon heating of the sample, Cu3Sn (peaks g and Z) and CuSe (peak h) sequentially

form at ~180 and ~190 C, respectively, via outward diffusion of Cu from Cu6Sn5, while

the Se layer crystallises. As soon as peaks g (Cu3Sn) with faint reflections at ~37° and ~41°

appear at ~180 C, peaks f (Cu6Sn5) weaken. Despite the weakening of peaks f, the

intensity of peak Z does not change along because the peak Z is also produced by β’-CuZn

located on the lower side of film. Furthermore, Cu3Sn which has the strongest reflection

near ~43° (peak Z) forms at ~180 °C. Therefore peak Z does not strengthen nor weaken

while peaks f and g change. After a while, peak h (CuSe) gradually emerges at ~190 C

along with peaks c (Sn) while peak f steadily disappears. This reaction clearly describes

the movement of Cu from Cu6Sn5 to the Se layer through the Sn layer (Cu6Sn5 → Cu3Sn

→ CuSe) due to the stacking order of the precursor for sample #3-6 (Mo/Cu/Zn/Cu/Sn/Se).

For this reason, the remaining Sn (peaks c) from the decomposition of Cu6Sn5 is observed

for a while. This movement of Cu is also in accordance with the reaction path for sample

#2-1 (Mo/Cu/Sn/Se), which has the same stacking order as that for the upper part of

sample #3-6. Here, CuSe formation at ~190 C at a temperature lower than the melting

point of Se indicates that this phase consists of crystalline Se and diffusing Cu, ie, Cu

reacts with the Se structure. Therefore, peak h may strengthen together with the detectable

Se reflections and reaches highest intensity as the Se reflections weaken at ~210 C.


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Figure 4.28: Time–temperature evolution of powder diffractograms of sample #3-6 (Mo/Cu/Zn/Cu/Sn/Se) with colour-coded intensities (a.u.) of the Bragg reflections. Phases denoted by peaks a-Z are as follows: d: SnSe, f: Cu6Sn5, g: Cu3Sn, h: CuSe, j: CuSe2, k: Cu2SnSe3, l: Cu2Se, n: ε-brass, K: Cu2SnSe3 (k), ZnSe and Cu2ZnSnSe4 (o), o: Cu2ZnSnSe4, Z: Cu, CuZn, Cu6Sn5 (b) and Cu3Sn (g).

At ~220 °C, CuSe2 and SnSe simultaneously form on the upper part of the film as

soon as Cu3Sn decomposes. The transformation of CuSe into CuSe2 starts at ~220 °C after

Se melts because of the increase in the amount of the more reactive Se (liquid). At the

same time, Cu diffused from Cu3Sn also forms CuSe2 with liquid Se. Because of the

consumption of Cu from Cu3Sn in CuSe2 formation, Cu3Sn decomposes. Consequently,

Sn decomposed from Cu3Sn also may react with liquid Se, forming the SnSe phase. In

comparison with those of the reaction path and the result for sample #2-1, the

decomposition temperatures for Cu6Sn5 and Cu3Sn in sample #3-6 are lower than those in

sample #2-1 probably because of the different amounts of Cu in Cu–Sn alloy. Because

around half of the Cu is used to form a Cu–Zn alloy in sample #3-6, another half of Cu is

consumed in the formation of Cu–Sn alloy. Contrary to this, the amount of Cu in Cu–Sn

alloy of sample #2-1 is double than the amount of Sn. A lower amount of Cu is included

in the Cu–Sn alloy of sample #3-6 than of sample #2-1. Therefore, the decomposition of

Cu3Sn, which signifies the formation temperature of SnSe, can occur in sample #3-6 at a

temperature lower than that for sample #2-1. It signifies that the Cu concentration in the

Cu–Sn alloy can influence on the formation temperature of SnSe.

Meanwhile, on the lower part of film, the dezincification progressively occurs in the

Cu–Zn alloy and leads the formation of ε-brass (Cu0.7Zn2) at ~250 C. Now peak Z above

~220 °C entirely belongs to β’-CuZn, because all Cu–Sn alloys (Cu6Sn5 and Cu3Sn)

decompose at ~220 C. Furthermore, the Sn–Zn–Cu phase diagram confirms that only β’-


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CuZn phase exists in the film among the Cu–Zn alloys near this temperature [17]. When

Cu–Sn alloy forms in the sample which consists of Cu, Zn and Sn elements, all Cu–Zn

alloys transform into the β’-CuZn phase, as described in section 2.1.1.4. This β’-CuZn

undergoes dezincification together with the outward diffusion of Cu as the temperature

rises. However, Zn interrupts Cu diffusion through the film, effectively stopping it,

trapping Cu in the Cu–Zn alloy and forming CurichZn phase on the upper side of Cu–Zn

alloy. While dezincification occurs, the dealloyed Zn forcedly moves to the bottom of the

Cu–Zn alloy because of outward diffusion of Cu (see section 4.1.3.4). Consequently, the

Cu concentration on the upper part of β’-CuZn increases while the Zn concentration on

the lower part of β’-CuZn increases. This sinking of Zn to the bottom of the β’-CuZn alloy

leads to the formation of ε-brass from ~250 C. Thus, peak n may be observed in Figure

4.28, similar to the result for sample #2-4a. Because ε-brass forms from β’-CuZn by

dezincification, strengthening of peak n (ε-brass) is inversely proportional to the

weakening of peak Z (β’-CuZn). This reaction path is also similar to that for sample #2-4a.

At ~270 C, CTSe forms from CuSe2 and SnSe slowly on the upper part of the film,

while ε-brass grows on the lower part of the film. Comparison with two results for sample

#2-1 (Mo/Cu/Sn/Se) and #2-4a (Mo/Cu/Zn/Se) suggests that the reflection at ~27° (peak

K) denotes only CTSe although peak K can also be produced by ZnSe. According to the

result for sample #2-1, CTSe can form from CuSe2 and SnSe at low formation rate, as

described in section 4.1.3.9. Similar to this formation reaction, Figure 4.28 indicates the

consumption of CuSe2 and SnSe phases at ~270 °C and above. When peak K emerges and

grows, peaks d (SnSe) and j (CuSe2) weaken in a manner inversely proportional to the

strengthening of peak K. This means that SnSe and CuSe2 form the phase, which is related

to peak K. On the contrary, peaks n and Z at around ~270 °C in Figure 4.28 do not match

well with the conditions for ZnSe formation which is found in the result for sample #2-4a.

According to the result for sample #2-4a, the conditions are as follows: i) achievement of

the lowest Bragg angle by peak n and ii) disappearance of peak Z as ZnSe forms via the

reaction of Se with Zn decomposed from the Cu–Zn alloy. At ~270 C, peak Z in Figure

4.28 is still detectable and peak n does not get the lowest Bragg angle but is moving to the

low angle side. Therefore, the beginning of peak K is clearly not due to ZnSe but rather

due to CTSe, following the reaction path for sample #2-1 (marked by ‘k’ to indicate

CTSe). One notable observation is the difference in formation temperatures of CTSe

(~270 C in sample #3-6 and ~290 C in sample #2-1). The reason of the difference

between two samples seems to be induced by the presence or absence of CuSe in the film.

As mentioned above, the Cu–Sn alloy in sample #3-6 has an amount of Cu lower than the

Cu–Sn alloy in sample #2-1 has. It is also revealed in the result for sample #2-1 that the

decomposition of this Cu–Sn alloy induces the formation of Cu–Se alloy. Thus, the CuSe

reflections (peaks h) are not detected in Figure 4.28 (sample #3-6) during CTSe formation,


85

whereas CuSe reflections are detected in Figure 4.4 (sample #2-1) because of the

difference in amounts of Cu in the Cu–Sn alloy between two samples. It suggests that the

absence of CuSe accelerates the reaction between CuSe2 and SnSe because CuSe2

facilitates the Se diffusion through film, as described in section 4.1.3.6.

At ~300 C, ZnSe gradually forms via decomposition of Cu–Zn alloy (peak Z). As

mentioned in the preceding paragraph, the formation temperature of ZnSe may be

confirmed by two conditions: the disappearance of peak Z and the lowest Bragg angle of

peak n because the shift of peak n to the lower Bragg angles indictates the increase in Zn

concentration in ε-brass. These two phenomena are observed at ~300 C in Figure 4.28.

After peak Z diminishes and vanishes at ~300 °C, peak n reaches the lowest Bragg angle

and then reverts to higher Bragg angles at ~330 °C. Because Zn from Cu–Zn alloy is

consumed in ZnSe formation from upper to lower part of Cu–Zn alloy, the Zn

concentration in CurichZn and ε-brass alloys sequentially decreases. Consequently, peak K

grows faster because of ZnSe formation.

At ~330 C, the formation rate of CTSe (peak k) increases as soon as CuSe2 (peaks j)

decomposes into CuSe (peak h). Consequently CuSe (peak h) and SnSe (peak d) disappear

at ~340 and ~350 C, respectively. When CuSe2 undergoes peritectic decomposition at

332 C [18], the weekening peaks j completely disappear, and one of the faint peaks h

(CuSe) appears on the right side of peak K as a shoulder peak of K. This faint reflection of

CuSe (peak h) is substantially weak in comparison with the other results, as observed

from several reflections of CuSe in the in situ XRD diffractograms, as in Figure 4.14

(sample #3-1) or 4.17 (sample #3-2). This faint reflection of CuSe shows that CuSe is

directly used to form CTSe as it dealloys from CuSe2. As it is revealed in section 4.1.3.9,

the rate of CTS formation varies with the Cu–Se alloys, and as it is known, the peritectic

decomposition of CuSe2 occurs at 332 C [18]. Thus, the rate of CTSe formation naturally

increases at ~330 C as the temperature rises, regardless of Cu concentration in the Cu–Se

alloy. That is, the reaction of SnSe with CuSe is faster than with CuSe2. As soon as CTSe

formation hastens, the requirement for CuSe increases. Therefore, CuSe is directly used to

form CTSe as it forms from the decomposition of CuSe2. In the meantime, peak d (SnSe)

weakens faster along with peaks j (CuSe2) and h (CuSe) and disappears at ~350 C soon

after peaks j and h disappear.

While ZnSe is continuously forms, ε-brass decomposes at ~360 °C, causing the

formation of Cu2Se. On the basis of the result for sample #2-4a (Figure 4.11), Cu2Se

forms from the decomposition of ε-brass. Likewise, the peaks l (Cu2Se) in Figure 4.28

also appears as soon as peak n (ε-brass) vanishes at ~360 °C. After ZnSe forms from

CurichZn on the upper side of Cu–Zn alloy, the Cu–Zn alloy has decomposed from the

upper to lower parts. Therefore ε-brass is resolved into Cu2Se and ZnSe at higher

temperature than vanishing temperature of peak Z (β’-CuZn). Meanwhile, faint peak h


86

(CuSe) slightly appears again on the high-angle side of peak K (~27°) at around 350–

370 C (Figure 4.28). This slight change implies that lots amount of Se diffuses into the

Cu–Zn alloy in accordance with the Cu–Se phase diagram [18], leading to the formations

of CuSe and Cu2Se. Afterwards, this CuSe directly converts into Cu2Se because of the

abundance of combinable Cu from the decomposed Cu–Zn alloy. Therefore Cu2Se may

form at ~360 °C although the temperature does not reach ~380 °C, which is near the

peritectic decomposition temperature of CuSe.

After Cu2Se forms at ~360 °C, it does not react with other components but steadily

grows in structural size. As shown in Figure 4.28, peak l (Cu2Se) does not weaken or

disappears since its appearance and remains until the end of the measurement.

Furthermore, peak l gradually shifts to lower Bragg angles as the temperature rises. This

indicates that Cu2Se does not combine with other components (CTSe and/or ZnSe), but

rather increases its unit cell size. According to equation (12) in section 4.2.1.5, the

absence of SnSe or SnSe2 may be a reason for the remaining Cu2Se in the film because

Cu2Se can combine with ZnSe together with SnSe or SnSe2. However all of SnSe

compounded in the sample #3-6 is already consumed for the CTSe formation at ~350 C –

peak d (SnSe) in Figure 4.28 disapppears at ~350 °C. This result clearly shows that Cu2Se

cannot combine with the CTSe and/or ZnSe phases, in accordance with the equation (12).

In other words, Cu2Se cannot form CZTSe without SnSe or SnSe2.

At ~380 C, CZTSe forms while the structure of Cu2Se grows. The emergence of

peaks o at ~35° and ~36° (Figure 4.28) verifies the formation of CZTSe at ~380 C. In

contrast to other results for samples #3-1 or #3-3, CZTSe in sample #3-6 forms not

immediately after the disappearance of SnSe but soon after the decomposition of ε-brass.


Figure 4.29: An arrow diagram for the reaction pathway of a Mo/CuZn/Cu/Sn/Se (sample #3-6) stacked layer during heat treatment. The rectangles with gradient color denote the decreasing (faint color) reflections of Se phase.


87

with gradient red colors indicates that the Se reflection decreases rapidly at ~210 °C and

disappear completely at ~220 °C. The arrows with dotted line from CuSe2 to an arrow for

Cu2SnSe3 through CuSe denote the reaction of small amount of CuSe. Because CuSe is

used to form CTSe immediately after the formation of CuSe from the decomposition of

CuSe2, only faint reflection of CuSe was observed in Figure 4.28. The letter with grey

color for ‘CuSe’ behind the letter for ‘Cu2Se’ expresses a short appearance of CuSe from

the decomposition of ε-brass. This CuSe soon transforms into Cu2Se.

4.2.3.3 Discussion

Two reaction paths of sample #3-5 (Mo/Zn/Cu/Sn/Se/Cu) and #3-6 (Mo/Cu/Zn/Cu/

Sn/Se) present the remaining Cu2Se in the synthesised film. In fact, the Cu2Se is well

known as a key component for the CuInSe2 formation as following equation [74]: Cu2Se +

In2Se3 → 2 CuInSe2. CZTSe comes from this CuInSe2 phase to replace indium (In) by

zinc (Zn) and tin (Sn). Therefore one can expect that the reaction process of CZTSe is

similar to that of CuInSe2. However, in contrast to CuInSe2, Cu2Se does not seem to be a

key reactant in the synthesis of CZTSe, but rather an interfering agent, in accordance with

the results for samples #3-5 and #3-6. As shown in Figures 4.26 (sample #3-5) and 4.28

(sample #3-6), the Cu2Se reflection does not disappear but remains detectable until the

end of the measurements, in contrast to the other previous results for sample #3-1–#3-4.

After the preparation of these two samples by division of the two Cu layers, β’-CuZn

is compounded thoroughly with those precursors as it is intended to delay the ZnSe

formation. Cu6Sn5 in each precursor is also detected. In particular, the transformation of

Cu6Sn5 into Cu3Sn by outward diffusion of Cu is also observed in the two samples.

According to the result for sample #2-1 (Mo/Cu/Sn/Se), this Cu diffusion through the Sn

layer leads to CTSe formation by forming CuSe2 and SnSe. On the basis of this formation

reaction of CTSe, two samples can also be expected to form CTSe since Cu3Sn is detected

in both samples. On the upper part of the Mo/Cu/Zn/Cu/Sn/Se film in sample #3-6, the

Cu/Sn/Se elemental layers follow a similar reaction path as that for sample #2-1. Thus

CTSe forms from CuSe2 and SnSe at ~270 C. Because of the lower concentration of Cu

in the Cu–Sn alloy, CTSe in sample #3-6 forms at a temperature lower than that required

for sample #2-1. In contrast to this, the Cu3Sn phase formed in sample #3-5 does not

decompose until ~410 C and transforms into Cu41Sn11 instead of CuSe2 and SnSe alloys.

The reason for the resistance to decomposition of Cu3Sn in sapmle #3-5 is the formation

of Cu-rich Cu–Se alloys (Cu2–xSe or Cu2Se phases) on the upper part of film, in particular

on the Cu3Sn alloy (Mo/alloys/Cu3Sn/Cu2–xSe), before Cu3Sn forms in the film.

According to the reaction path of sample #3-5, Cu3Sn remains under Cu–Se alloys

because Cu cannot easily diffuse into Cu2–xSe alloys. Furthermore, the reaction path of


88

this sample at ~340 °C shows that, as soon as Cu2Se forms, Cu diffusion through the Cu–

Se alloy becomes entirely blocked, leading to slightly strengthened Cu3Sn reflections. At

the same temperature, Se diffusion into the film is also interrupted as the Cu2Se phase

stabilises. The reason for the interruption of Se diffusion seems to be the absence of CuSe2

in sample #3-5 because it facilitates the Se diffusion into the film (see section 4.1.3.5).

Therefore none of the metallic elements can react with Se. Only certain amount of ZnSe

forms while Cu2–xSe transforms into Cu2Se. The unchanged intensity of ZnSe reflections

until the end of the measurement also verifies the cessation of Se diffusion through film

(see Figure 4.26). For this reason, Cu3Sn cannot react with Se to form Cu–Se and Sn–Se

alloys and CTSe thus cannot form in sample #3-5, in contrast to sample #3-6. One

conclusion that can be derived from this result is that Cu2Se interrupts Se diffusion

through the film.

According to the formation reaction of Cu2Se in sample #3-5, Cu2Se derives from the

Cu3Se2 phase compounded in initial precursor through the following reaction: Cu3Se2 →

Cu2–xSe → Cu2Se. As Cu3Se2 decomposes at 113 C into Cu2–xSe [18], Cu2Se may form

at low temperature before other binary selenides form in the film. This signifies that the

formation of Cu3Se2 in the precursor drives the earlier formation of Cu2Se in the film and

eventually ceases the Se diffusion through the film. In contrast to sample #3-5, sample #3-

6 has only amorphous Se layer in the initial precursor, and subsequently CuSe2 forms via

CuSe formation on the upper side of film. As revealed in section 4.1.3.5, CuSe2 facilitates

Se diffusion. Consequently, Cu3Sn and Cu–Zn alloy phases in sample #3-6 can react with

Se and form CuSe2 and SnSe from Cu3Sn and ZnSe from Cu–Zn alloy in the film.

Therefore CZTSe can eventually form, in contrast to sample #3-5. Comparison the

different results between these two samples reveals that the CuxSey phase formed at low

temperature in the film can influence CZTSe formation. Thus preventing the formation of

Cu2Se and/or Cu3Se2 phases in the film at low temperature is necessary.

The unfriendly incompatibility of Cu2Se with other components without SnSe or

SnSe2 is also observed in sample #3-6 during formation of CZTSe. The reaction path of

sample #3-6 on the bottom of the film suggests that β’-CuZn mainly forms in the initial

precursor and converts into ε-brass by dezincification and outward diffusion of Cu. As Se

reacts with Zn of the Cu–Zn alloy, the Cu concentration in the alloy steadily increases.

Therefore, Cu2Se easily forms because of the increased concentration of Cu in the Cu–Zn

alloy when Cu in the Cu–Zn alloy can also react with Se. After Cu2Se forms, it does not

combine with other components such as ZnSe or CTSe, but its structure grows until the

end of measurement. Equations (11) and (12) (see section 4.2.1.5) for the formation of

CZTSe indicate that Cu2Se can react with ZnSe only in the presence of SnSe (or SnSe2),

whereas CTSe reacts with ZnSe, ie, Cu2Se needs SnSe (or SnSe2) for the reaction.

According to the reaction path for sample #3-6, however, SnSe disappears at ~350 °C


89

before Cu2Se appears at ~360 C. This means that when Cu2Se forms in the film, no SnSe

for the reaction of Cu2Se is left; thus, Cu2Se remains inside the film. It seems that the

separate formation of the Cu–Sn alloy from the Cu–Zn alloy in the precursor induces the

precipitation of Cu2Se by reaction with all of the SnSe for CTSe formation. The reaction

path of sample #3-3 clearly shows the consumption of Cu2Se for the CZTSe formation by

SnSe (see Figure 4.19). Although Cu–Zn alloy in sample #3-3 undergoes decomposition

into the ZnSe and Cu2Se phases at ~350 °C, similar to the formation of Cu2Se in sample

#3-6, Cu2Se in sample #3-3 combines with other components whereas Cu2Se in sample

#3-6 remains in the film. The reason of this different result for sample #3-3 is the

simultaneous formation of SnSe and Cu2Se. Because sample #3-3 forms the eutectic Sn–

Zn alloy (instead of Cu–Sn alloy) together with Cu–Zn alloy in the film, SnSe forms at

temperature higher than that for sample #3-6. Therefore the SnSe can be used to combine

with this Cu2Se. The comparison of formation of Cu2Se between samples #3-3 and #3-6

presents that separate preparation of Cu–Sn and Cu–Zn alloys in the precursor is counter-

intuitive.

In conclusion, Cu2Se interrupts Se diffusion through the film and reacts with ZnSe

only when SnSe (or SnSe2) is near these two components. Consequently, the Cu2Se (and

Cu2–xSe) phase must not be formed before other binary selenides are formed in the film.

Thus, formation of Cu3Se2 in the initial precursor should be avoided.


90

4.3 Secondary phases in Cu2ZnSnSe4 films: How and why they form

In this section, the formation reactions of each component for CZTSe is discussed on

the basis of the reaction paths investigated in sections 4.1 and 4.2. In particular, the cause

of the remaining secondary phases in the film is explained by an understanding of these

reactions.

4.3.1 Cu2Se

It is well known that the Cu2Se is formed at 379.7 °C by decomposition of CuSe [18].

This formation reaction is obviously observed in the result for sample #2-3

(Mo/Zn/Cu/Se). However, three-metal system does not show this reaction path because

lots amount of CuSe is consumed for CTSe or CZTSe formation before it decomposes

into Cu2Se. The Cu2Se in three-metal system is generally formed below ~380 °C by the

decomposition of Cu–Zn alloy, in accordance with the results for sample #3-3 and #3-6.

Only when Cu3Se2 is compounded in the precursor, as in the case of sample #3-5, Cu2–xSe

may form at 113 °C [18], causing the immiscible Cu2Se with other components.

The main reason of residual Cu2Se in the film is the absence of Sn–Se alloy near

Cu2Se. Equation (12) in section 4.2.1.5 implies that Cu2Se can combine with ZnSe only in

the presence of SnSe (or SnSe2) to form CZTSe. It cannot combine with ZnSe only

because the Cu–Zn–Se alloy does not form in nature. Additionally, SnSe2 may form from

SnSe under high pressure. This indicates that the existence of the SnSe alloy near Cu2Se

and ZnSe phases is needed to remove the remaining Cu2Se from the CZTSe film. The

reaction paths for samples #3-5 and #3-6, which have a residual Cu2Se in the synthesised

film, suggest that the cause of remaining Cu2Se is the absence of SnSe alloy near Cu2Se.

The earlier formation of Cu2–xSe in sample #3-5 (Mo/Zn/Cu/Sn/Se/Cu) also induces

the absence of SnSe near Cu2Se in the film by disturbing Se diffusion through the film. As

shown in section 4.2.3.1 (sample #3-5), Cu2Se comes from Cu2–xSe, which is formed by

decomposition of Cu3Se2 at 113 C. As Cu2–xSe forms in the film, Se diffusion through

film becomes restrictive. Thus, only certain amount of Se may react with Zn when Cu2–

xSe decomposes into Cu2Se. Although Se diffusing from Cu2–xSe is locationally more

closed to Cu3Sn than Cu–Zn alloy, the Se reacts with Zn instead of liquid or gaseous Sn

because of the reaction tendency (Sn–Se < Zn–Se, see section 4.1.3.10). That is, earlier

formation of Cu2–xSe before SnSe formation causes Cu2Se to remain in the film.

The SnSe formation during the presence of CuSe2 phase in the film seems to

consequently prevent the combination of Cu2Se with SnSe, in accordance with sample #3-

6 (Mo/Cu/Zn/Cu/Sn/Se). Because SnSe easily react with CuSe2 to form CTSe, all of SnSe


91

in sample #3-6 is already consumed before the Cu2Se is formed by decomposition of Cu–

Zn alloy at ~360 °C. Because of the use of all of SnSe before Cu2Se formation, the Cu2Se

may not combine with other components but remains in the film. Contrary to Cu2Se in

sample #3-6, the Cu2Se in sample #3-3 (Mo/Zn/Sn/Cu/Se) is observed at 350–360 °C and

disappears. Because SnSe and Cu2Se simultaneously form at ~350 °C in sample #3-3, the

Cu2Se may immediately combine with SnSe for the CTSe or CZTSe formation. The

difference in results between samples #3-3 and #3-6 proves again that the absence of SnSe

near Cu2Se induces the residual Cu2Se in the film.

In conclusion, the earlier formation of Cu2–xSe (or Cu3Se2 formation in the precursor)

before the formation of Sn–Se alloy and the formation of SnSe together with CuSe2

(instead of Cu2Se) leave Cu2Se in the film. Although Cu2Se is a known super-ionic

conductor for the synthesis of chalcopyrite [74] which is chemically identical to kesterite,

it is not used in the synthesis of kesterite structure unless SnSe (or SnSe2) and ZnSe alloys

are compounded near Cu2Se in the film. In fact, the decomposition of CZTSe is also a

reason for the residual Cu2Se in kesterite film [56], but this is not discussed in this section

because this study examines only the formation reaction.

4.3.2 ZnSe

ZnSe is an alloy with the easiest formation because Zn has strongest affinity among

the metallic elements to combine with Se. Therefore, if Se is in contact with three metals

simultaneously during annealing, Se reacts in sequence with Zn, Cu and then Sn, in

accordance with the Ellingham diagram and the conclusion in section 4.1.3.10 (Sn–Se <

Cu–Se < Zn–Se). Therefore, when the metal–Zn alloy decomposes during Se diffusion

through film, ZnSe forms before Se reacts with another metal element. For example, when

the Cu–Zn alloy decomposes, Cu2Se may form after the formation of ZnSe phase, in

accordance with the result for sample #2-4 (Mo/Cu/Zn/Se). When there is no obstacle for

the reaction between Zn and Se, as in sample #1-1 (Mo/Zn/Se), ZnSe can form at ~290 °C.

Since ZnSe easily forms earlier than other binary selenides, it is slightly difficult to

combine with other components for CZTSe formation as ZnSe structure becomes a stable

state before other components form (see section 4.2.1). Therefore it is necessary to delay

the ZnSe formation more than the formation of the other components for CZTSe

formation. According to equations (11) and (12) in section 4.2.1.5, the combinable

components for CZTSe with ZnSe are CTSe or Cu2Se and SnSe (or SnSe2). However,

only the formation temperature of CTSe is worth considering because Cu2Se, SnSe and

SnSe2 generally form later than ZnSe. The Cu2Se is well-known to forms at 379.3 C via

decomposition of CuSe unless the proportion of Se in Cu–Se alloy is less than ~40 at%

[18]. The SnSe phase typically forms after ZnSe formation due to the strong affinity of Se


92

to Zn (Sn–Se < Zn–Se), as revealed in section 4.1.3.10. SnSe2 may be dismissed because

this phase forms after SnSe formation under high pressure. Contrary to these three phases,

CTSe can form at ~290 °C regardless of stacking order of Cu and Sn layers in the

precursor, in accordance with the results for samples #2-1 (Mo/Cu/Sn/Se) and #2-2

(Mo/Sn/Cu/Se). This formation temperature of CTSe is the same as that of ZnSe from the

result for sample #1-1. Thus, it is necessary to delay the ZnSe formation more than CTSe

formation for the reduction or absence of residual ZnSe in the film. Additionally, the

formation temperatures of these two phases can be changed with the stacking order of

elemental layers in the initial precursor.

One approach to delaying ZnSe formation is the formation of Cu–Zn alloy in the

initial precursor. The sample #3-3 and #3-6, which have a β’-CuZn phase in the precursor,

proves the ZnSe formation later than the CTSe formation. As temperature rises, the Cu–

Zn alloy in these two samples produces Cu-rich Cu–Zn alloy (such as Cu2Zn) under

CuSe2 by dezincification. Because Cu in the Cu–Zn alloy traps and isolates Zn from the

reaction with Se, as the amount of Cu in the Cu–Zn alloy near CuSe2 increases, Zn needs

more thermal energy to react with Se diffusing from CuSe2. Thus, the formation

temperature of ZnSe is changed depending on the Cu concentration in the Cu–Zn alloy

near Se. Accordingly, sample #2-4a (Mo/Cu/Zn/Se, [Cu]/[Zn] = 1.3) forms ZnSe at

~300 °C whereas sample #2-4 (Mo/Cu/Zn/Se, [Cu]/[Zn] = 2) forms it at ~360 °C. On the

other hand, the formation of Cu–Zn alloy in the precursor can also induce the residual

Cu2Se in the film. According to the result for sample #3-6, the Cu2Se cannot be consumed

for CZTSe formation bur remains in the film. The reason of it was the absence of eutectic

Sn–Zn alloy near Cu–Zn alloy. Therefore, the formation of Cu–Zn alloy near the Sn–Zn

mixture during heating of the sample is necessary to delay the ZnSe formation, as well as

for the absence of Cu2Se in the film, as described in section 4.2.1.5.

On the other hand, separation between Zn and Se layers by Cu or Sn layers may be

another way to delay the ZnSe formation. However, that separation does not seem to

accomplish the delay every time. On the Cu layer for the dividing Zn and Se layers, such

as that in sample #2-3 (Mo/Zn/Cu/Se), ZnSe forms at ~290 C because of the facile

diffusion of Se through CuSe2 (see section 4.1.3.5). Sample #3-2 (Mo/Zn/Sn/Se/Cu),

which deposits Sn layer in between Zn and Se layers, also forms ZnSe at ~290 °C because

the Sn and Zn layers form a eutectic Sn–Zn alloy with different distribution of Zn through

this Sn–Zn alloy. On the contrary, sample #2-6 (Mo/Zn/Sn/Se), which also prepares Sn

layer in between Zn and Se layers, exhibits ZnSe formation at ~350 °C. Depending on the

liquidus temperature of Sn–Zn alloy which is determined by the concentration of Sn in its

alloy, the formation temperature of ZnSe is determined.

In conclusion, it is necessary to form ZnSe after the formation of other components of

CZTSe, especially CTSe. Moreover, only the formation of Cu–Zn alloy in the precursor


93

delays ZnSe formation, but caution must be taken for the residual Cu2Se by absence of

Sn–Zn alloy in the film.

4.3.3 SnSe2 (or SnSe)

The reaction of Se with Sn can occur after Se reacts with Zn and Cu, in accordance

with the Ellingham diagram and with the tendency of the Se reaction (Sn–Se < Cu–Se <

Zn–Se, see section 4.1.3.10). Although Se can react with liquid Sn at ~230 °C, in

accordance with the result for sample #1-2 (Mo/Sn/Se), SnSe appears always after the

formation of ZnSe and Cu–Se alloy phases. In general, SnSe is formed by decomposition

of metal–Sn alloy. As the Cu–Sn alloy decomposes due to the outward diffusion of Cu,

SnSe may form after the formation of Cu–Se alloy. Similar to this, SnSe also forms after

the ZnSe formation as the eutectic Sn–Zn alloy melts.

Another Sn–Se alloy, SnSe2, forms from SnSe but is easily affected by sample

pressure. The result for sample #1-2 clearly shows an influence of the pressure on the

formation of Sn–Se alloy. When the sample bursts at ~520 °C, the SnSe2 which has

formed from SnSe at ~270 °C reconverts into SnSe (see Figure 4.2 in section 4.1.1.2).

Because of the sudden decrease in sample pressure by its rupture, SnSe2 cannot maintain

its phase but transforms into SnSe again. It signifies that SnSe2 may form only under high

pressure.

Variation of the formation of Sn–Se alloys with pressure seems to influence the

residual SnSe or SnSe2 phase in the kesterite film. It seems that Cu2Se prefers to react

with SnSe2 rather than SnSe. When some measurements are performed under low pressure

because of a loosely clamped sample holder, the SnSe reflection does not disappear until

the end of measurement. The peak d (SnSe) observed in the preceding results for three-

metal system generally vanishes as distinguishable reflections of CZTSe (peaks o) emerge.

For example, peak d in Figure 4.19 (sample #3-3: Mo/Zn/Sn/Cu/Se) also disappears as

peaks o (CZTSe) appear. On the contrary, peak d in Figure S1 diminishes and increases

again in spite of the appearance of peaks o (CZTSe) in the diffractograms. This Figure S1

described in section for supplementary information is another result for sample #3-3

which is annealed under low pressure. Accordingly, the sudden shift of all reflections at

the same temperature, which indicates the rupture of sample, is not observed in Figure S1,

in contrast to Figure 4.19, because the evaporating Se gas from sample may easily leak.

The difference in results between two measurements implies that the low sample pressure

causes the residual SnSe in the kesterite film. Considering the decomposition of CuSe into

Cu2Se at 379.7 °C [18] and the transformation of SnSe2 into SnSe under low pressure, it

seems that Cu2Se is reluctant to react with SnSe. For this reason, SnSe seems to remain in

the synthesised film only when the sample is annealed under low pressure. Although the

transformation of CuSe into Cu2Se is not detected in Figure 4.19, it seems that a certain

amount of CuSe transforms into Cu2Se as a nano-crystalline structure: the Cu2Se

reflection in Figures 4.19 and 4.28 is due to the decomposition of Cu–Zn alloy and not

because of the decomposition of CuSe. Furthermore, the peak d may denote not only SnSe

but also SnSe2 because the main reflections of SnSe and SnSe2 have the same Bragg angle


94

at ~30°. Thus unobservable reflections of Cu2Se may also signify the reaction between

Cu2Se, SnSe2 and ZnSe as soon as CuSe converts into Cu2Se. Therefore, the sample

pressure needs to be high enough as much as SnSe may transforms into SnSe2. The need

of high pressure during annealing is in accordance with other literature [5] which reveals

the evaporation of SnSe during heating of the sample. The detectable SnSe reflection in

Figure S1 differs from this investigation [5], but the necessity of high pressure concurs

with it.

According to other studies [75, 76], Sn(S,Se)2 is sometimes observed in the kesterite

film as a secondary phase. The reason of this remaining SnSe2 in the film is not obviously

revealed in this study. However, the SnSe2 seems to occur by decomposition of Cu2SnSe3

into Cu4SnSe4 [5] under high pressure. On a basis of other study [5] which is mentioned

above, SnS may evaporate from the film above 350 °C as Cu2SnS3 transforms into

Cu4SnS4. That means when CTSe cannot react with ZnSe for some reason, such as the

stabilized structure of ZnSe before the formation of CTSe (see section 4.2.1), a part of

CTSe would decompose into Cu4SnSe4, leading to the separation of SnSe phase. This

SnSe seems to convert into SnSe2 due to the high pressure during the synthesis of CZTSe,

causing the residual SnSe2 in the film.

The Cu-poor and Zn-rich composition of CZTSe film seems to also be another reason

for the remaining SnSe or SnSe2 phases in the kesterite film. Considering the elemental

ratio of three-metal samples, which also consists of Cu-poor and Zn-rich compositions for

this study, the residual SnSe or SnSe2 phase in the kesterite film seems to be inevitable as

following equation (13):

18 Cu + 12 Zn + 10 Sn + 53 Se → 9 Cu2ZnSnSe4 + 3 ZnSe + SnSe/SnSe2 + 13/12 Se (13)

However, the SnSe or SnSe2 is not observed at the end of measurements for three-metal

samples, as observed in section 4.2. Only when the sample anneals under low pressure,

SnSe or SnSe2 is detected together with CZTSe (see Figure S1).

In conclusion, high pressure is necessary to transform the SnSe into SnSe2 for the

consumption of Sn–Se alloy into CZTSe due to the preference of Cu2Se for its reaction.

Additionally, the non combination of CTSe and ZnSe before the decomposition of CTSe

into Cu4SnSe4 and/or the Cu-poor and Zn-rich compositions of sample seems to induce

the residual SnSe or SnSe2 in the kesterite film.

4.3.4 Cu2SnSe3

On the basis of this study, CTSe can be form via reactions (r-i) between CuSe and

SnSe, (r-ii) between CuSe2 and SnSe, and (r-iii) between CuSe2 and liquid Sn. As shown

by results for samples #2-1 (Mo/Cu/Sn/Se) and #2-2 (Mo/Sn/Cu/Se), CTSe may form at

~290 C if there is no interference with reaction (r-ii). Additionally, the formation rate of

the reaction (r-ii) is slower than that of reaction (r-i) (see section 4.1.3.9). Only when


95

SnSe forms near the CuSe2 phase, as in the case of sample #3-6, the formation

temperature of CTSe reaches ~270 C which is lower than its common value (~290 °C) –

the SnSe observed in this study generally forms after the decomposition of CuSe2. This

signifies that if SnSe is already compounded in the precursor together with CuSe2 or CuSe,

then the formation temperature of CTSe can be much lower than ~290 C. Reaction (r-iii)

may occur only when the Cu–Zn and Sn–Zn alloys are compounded in the film, in

accordance with the result for sample #3-3. This previous result shows the reaction

between CuSe2 and liquid Sn at ~300 °C as soon as the eutectic Sn–Zn alloy liquefies.

Meanwhile, the Zn from the molten Sn–Zn alloy adheres to the Cu–Zn alloy because of

the formation of Cu2Zn under CuSe2 (see section 4.2.1.5).

The CTSe phase is rarely observed as a residual alloy in the kesterite film. Such

observation signifies the lack of Zn or insufficient annealing time because CZTSe usually

forms from CTSe and ZnSe phases and generally decomposes into Cu2Se, SnSe and ZnSe

[56].

4.3.5 Cu2ZnSnSe4

The data suggest that CZTSe forms from CTSe and ZnSe in general, consistent with

equation (11) written in section 4.1.2.5. CZTSe formation from Cu2Se, ZnSe and SnSe2

(or SnSe) components, which follows equation (12), is not obviously observed in this

study. Only the undetectable reflection of Cu2Se (peak l) and the diminishing and

vanishing reflection of SnSe (or SnSe2) (peak d) in the diffractograms for three-metal

system may infer the reaction following equation (12) above ~380 °C. It seems that the

nano-crystalline Cu2Se forms from the decomposition of CuSe at 379.7 °C [18] and

immediately reacts with SnSe (or SnSe2), leading to the unobservable peak l and the

disappearing peak d. After the peak d disappears, peaks o generally appear at 380–420 °C.

It indicates that the last reaction for the formation of CZTSe is the consumption of SnSe

which always forms as a last binary selenide (see section 4.3.3). Anyhow, the formation of

CZTSe below ~380 °C generally follows equation (12) unless Cu2–xSe (or Cu2Se), ZnSe

and SnSe2 (or SnSe) phases are prepared in the precursor.

To synthesise the pure CZTSe film, the formation of Cu–Zn and Sn–Zn alloys in the

film during annealing under high pressure is necessary to reduce or remove the residual

Cu2Se, ZnSe and SnSe phases in the kesterite film, as described in section 4.3.1–4.3.3. It

may done by preparation of Sn layer in between Cu and Zn layers, such as sample #3-3

(Mo/Zn/Sn/Cu/Se). Other study [77] also presents that the sequence of Mo/Cu/Sn/Zn

draws the 93% pure CZTSe film among other samples which have various sequences.

Therefore, it is valuable to synthesis the kesterite film, which consists of an exact

composition of CZTSe with a Mo/Zn/Sn/Cu/Se sequence, under high pressure.

5. Conclusion

96

5. Conclusions

To understand the mechanism of reaction between the elements Cu, Zn, Sn and Se,

various samples with different numbers of metallic layers and a Se layer, as well as

different stacking orders, are prepared. These samples are measured by time-resolved in

situ XRD to observe the reaction paths. As expected, reaction paths generally vary with

the sequence of stacking layers. Observation of the reaction paths for one- and two-metal

samples (section 4.1) reveals the tendencies for reaction between the four elements, as

follows:

Sn–Se < Sn–Zn < Cu–Sn < Cu–Se < Cu–Zn < Zn–Se

In fact, the stronger tendency of Sn to react with Cu than with Zn (Sn–Zn < Cu–Sn) is

not obviously revealed in this study. However, its ordering indicated above reflects the

fact that eutectic Sn–Zn alloy does not crystallise whereas Cu–Sn alloy does. The

tendency of Zn to react with Se than with Cu (Cu–Zn < Zn–Se) may be applicable

depending on the Cu concentration in the Cu–Zn alloy near Se element (see section

4.1.3.7). As the Cu concentration increases, the formation temperature of ZnSe formed by

the decomposition of Cu–Zn alloy also increases. On the basis of these reaction tendencies

of the four elements, results for the three-metal samples are analysed.

CZTSe in this study mostly forms from CTSe and ZnSe. In particular, the formation of

CTSe before the crystallisation of ZnSe phase is necessary to reduce or remove the

residual ZnSe in the kesterite film. ZnSe generally forms at ~290 °C without interferences.

CTSe also may form at ~290 °C from CuSe2 and SnSe. However the formation of SnSe is

the difficult reaction because Se has weakest tendency to react with Sn than with Zn or

with Cu (Sn–Se < Cu–Se < Zn–Se), as described in above. Therefore ZnSe generally

forms earlier than the CTSe forms, resulting in the remaining ZnSe. For this reason, it is

necessary to delay the ZnSe formation after or simultaneously with the formation of CTSe

(or Cu2Se and SnSe/SnSe2).

The way to delay the ZnSe formation is the formation of Cu–Zn alloy in the precursor.

However, the formation of Cu–Zn alloy separate from the Cu–Sn alloy should be avoided

because it may induce the residual Cu2Se in the film. The way to prevent the remaining

ZnSe and Cu2Se phases in the film is the formation of eutectic Sn–Zn alloy near the Cu–

Zn alloy. The co-existence of these two phases may draw by deposition of the Sn layer

between Cu and Zn layers without division of the Cu layer into two layers.

The residual Cu2Se in the film occurs because of absence of SnSe or SnSe2 near Cu2Se.

When Cu2Se or Cu2–xSe forms at an early stage of the reaction of CZTSe, the Se diffusion

through film is interrupted, preventing the SnSe formation. Consequently, Cu2Se remains

5. Conclusion

97

in the absence of CZTSe structure (section 4.2.3.1) because the Cu2Se phase can combine

with SnSe or SnSe2 together with ZnSe for the CZTSe formation. To prevent a formation

of residual Cu2Se, non formation of Cu3Se2 or Cu2–xSe phases or the formation of Sn–Se

alloy in the precursor is recommened.

The reason of residual SnSe or SnSe2 seems to be correlated with the sample pressure

during a synthesis of CZTSe. When the sample is annealed under low pressure, SnSe

remains in the kesterite film (section 4.3.3). Additionally, the SnSe2 may transform into

SnSe under low pressure (section 4.1.3.1). Considering the decomposition of CuSe into

Cu2Se at 379.7 °C, the Cu2Se seems to hesitate to react with SnSe rather than prefer to

react with SnSe2. Therefore the high pressure is necessary to convert SnSe into SnSe2 for

the integration of Sn–Se alloy into CZTSe. The reason for a detectable SnSe2 in the

kesterite film is uncertain. It seems to occur while CTSe decomposes into Cu4SnSe4 under

high pressure before CTSe reacts with ZnSe. Therefore the combination between CTSe

and ZnSe is necessary again.

In conclusion, the path of the reaction between the four elements and, thus, the

homogeneity of the kesterite film, is influenced by the sequence of stacking layers in the

initial precursor. According to this study, several precursors may be used to improve the

kesterite film, but only one sequence of Mo/Zn/Sn/Cu precursor is recommended here.

Because Cu may easily react with Se, the Cu layer on top of film may absorb lots amount

of Se more than other metallic layers. This recommended precursor leads to Cu2SnSe3

formation next to the ZnSe under Cu–Se alloy. Therefore, this sample may be used to

achive a homogeneous Cu2ZnSnSe4 film.

References

98

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Appendix

105

Appendix

Supplementary information

Figure S1: Time–temperature evolution of another powder diffractograms for sample #3-1 (Mo/Zn/Sn/Cu/Se) with colour-coded intensities (a.u.) of Bragg reflections. Phases denoted by peaks a-Z are as follows: c: Sn, d: SnSe, h: CuSe, j: CuSe2, k: Cu2SnSe3, o: Cu2ZnSnSe4, K: ZnSe, Cu2SnSe3, and Cu2ZnSnSe4 (f), Z: Cu, Zn, Cu5Zn8, CuZn, and α-brass (Cu2Zn). Because sample anneals under low pressure, the diminished SnSe appears again and is steadily detected by the end of this measurement.

Figure S1 is another result for sample #3-3 (Mo/Zn/Sn/Cu/Se) which is annealed

under low pressure due to the loosely clamped sample holder. Accordingly, the sudden

shift of all reflections at the same time, which obviously indicates the rupture of sample, is

not observed in Figure S1, in contrast to Figure 4.19 in section 4.2.1.3, because

evaporating Se and/or SnSe may easily leak from the sample. Therefore the interpretation

of in situ analysis for Figure S1 is the same as that for sample #3-3 described in section

4.2.1.3.

One difference in results between two measurements is the reaction path of SnSe. As

shown in Figure S1, the reflection of SnSe (peak d) diminishes and increases again along

with the reflections of CZTSe (peaks o and K). In contrast, peak d in Figure 4.19

diminishes and disappears as soon as peaks o emerge at ~400 °C. That means only when

the sample anneals under low pressure, the reflection of SnSe is detectable until the end of

measurement. In fact, this peak d may also denote SnSe2 because the main reflections of

SnSe and SnSe2 have the same Bragg angle at ~30°. Therefore the non-disappearance of

peak d in Figure S1 signifies that the Sn–Se alloy may not consume into CZTSe when the

synthesis of CZTSe are performed under low pressure.

Appendix

106

List of publications

As a First Author (4)

Hyesun Yoo, Arnaud Verger, Robert Lechner, Virginie Moreau, Stefan Jost, Jörg Palm,

Rainer Hock

“Different reaction pathway for the formation of Cu2ZnSnSe4 thin film from

different stacking order of elemental layers”

Proceedings of the 29th European Photovoltaics Solar Energy Conference and Exhibition

(EU-PVSEC 2014), pp.1477-1482, DOI: 10.4229/EUPVSEC20142014-3BO.7.3.

H. Yoo, R. Lechner, S. Jost, J. Palm, A. Verger, A. Lelarge, V. Moreau, C. Papret, R.

Hock

“The effect of secondary phases on Cu2ZnSn(S,Se)4 based solar cell”

Photovoltaic Specialist Conference (PVSC), 2014 40th IEEE, pp.2431-2435,

DOI: 10.1109/PVSC.2014.6925420.

H. Yoo, R.A. Wibowo, G. Manoharan, R. Lechner, S. Jost, A. Verger, J. Palm, R. Hock

“The formation mechanism of secondary phases in Cu2ZnSnSe4 absorber layer”

Thin Solid Films 582 (2015) 245-248, DOI: 10.1016/j.tsf.2014.08.048.

Hyesun Yoo, R.A. Wibowo, A. Hölzing, R. Lechner, J. Palm, S. Jost, M. Gowtham, F.

Sorin, B. Louis, R. Hock

“Investigation of the solid state reactions by time-resolved X-ray diffraction

while crystallizing kesterite Cu2ZnSnSe4 thin films”

Thin Solid Films 535 (2013) 73-77.

As a co-author (6)

Urike Künecke, Christina Hetzner, Stefan Möckel, Hyesun Yoo, Rainer Hock, Peter

Wellmann

“Characterization of kesterite thin films fabricated by rapid thermal processing of

stacked elemental layers using spatially resolved cathodoluminescence”

Thin Solid Films 582 (2015) 387-391.

Rachmat Adhi Wibowo, Stefan Möckel, Hyesun Yoo, Astrid Hölzing, Rainer Hock,

Peter J. Wellmann

“Formation of Cu2SnSe3 from stacked elemental layers investigated by combined in situ

X-ray diffraction and differential scaning calorimetry techniques”

Journal of Alloys and Compopunds 588 (2014) 254-258.

Appendix

107

Rachmat Adhi Wibowo, Stefan A. Möckel, Hyesun Yoo, Christina Hetzner, Astrid

Hölzing, Peter Wellmann, Rainer Hock

“Intermetallic compounds dynamic formation during annealing of stacked elemental

layers and its influences on the crystallization of Cu2ZnSnSe4 films”

Materials Chemistry and Physics 142 (2013) 311-317.

R.A. Wibowo, H. Yoo, A. Hölzing, R. Lechner, S. Jost, J. Palm, M. Gowtham, B. Louis,

R. Hock

“A study of kesterite Cu2ZnSn(Se,S)4 formation from sputtered Cu-Zn-Sn metal

precursors by rapid thermal processing sulfo-selenization of the metal thin films”


R. Lechner, S. Jost, J. Palm, M. Gowtham, F. Sorin, B. Louis, H. Yoo, R.A. Wibowo, R.

Hock

“Cu2ZnSn(S,Se)4 solar cells processed by rapid thermal processing of stacked elemental

layer precursors”


A. Hölzing, R. Schurr, H. Yoo, R.A. Wibowo, R. Lechner, J. Palm, S. Jost, R. Hock

“Real-time investigation on the formation of Cu(In,Ga)(S,Se)2 while annealing Cu-In-

Ga precursors with different sulphur-selenium mixtures”

Thin Solid Films 535 (2013) 112-117.

During the master thesis (5)

Hyesun Yoo, JunHo Kim, Lixin Zhang

“Sulfurization temperature effects on the growth of Cu2ZnSnS4 thin film”

Current Applied Physics 12 (2012) 1052-1057.

Hyesun Yoo, JunHo Kim

“Growth of Inx(S, O, OH)y films by chemical bath deposition”

Current Applied Physics 11 (2011) S81-S87.


“Comparative study of Cu2ZnSnS4 film growth”

Solar Energy Materials & Solar Cells 95 (2011) 239-244.


“Growth of Cu2ZnSnS4 films by sputtering with post-sulfurization”

AIP Conference Proceedings 157 (2011) 1399, DOI: 10.1063/1.3666304.

Appendix

108


“Growth of Cu2ZnSnS4 thin films using sulfurization of stacked metallic films”

Thin Solid Films 518 (2010) 6567-6572.

Dong-Yeup Lee, Hyesun Yoo, Ki-Bong Song, Jae Ho Yun, JunHo Kim

“Growth of sprayed CIS film and solar cell application”

Photovoltaic Specialists Conference (PVSC), 2010 35th IEEE, Honolulu, HI, pp.003443-

003445, DOI: 10.1109/PVSC.2010.5614636.

Conference contributions

Hyesun Yoo, Arnaud Verger, Robert Lechner, Virginie Moreau, Stefan Jost, Jörg Palm,

Rainer Hock

“Different reaction pathway for the formation of Cu2ZnSnSe4 thin film from different

stacking order of elemental layers”

» Best Student Award «

Talk, 29th European Photovoltaic Solar Energy Conference and Exhibition (EU PVSEC

2014), RAI Convention and Exhibition Centre, Amsterdam, September 22-26, 2014

H. Yoo, R. Lechner, S. Jost, J. Palm, A. Verger, A. Lelarge, V. Moreau, C. Papret, R.

Hock

“The effect of secondary phases on Cu2ZnSn(S,Se)4 based solar cell”

Poster, 40th IEEE Photovoltaic Specialists Conference (PVSC), Colorado Convention

Center, Denver, Colorado, June 8-13, 2014.

H. Yoo, R.A. Wibowo, G. Manoharan, R. Lechner, S. Jost, A. Verger, J. Palm, R. Hock

“The formation mechanism of secondary phases in Cu2ZnSnSe4 absorber layer”

Talk, E-MRS 2014 Spring Meeting, Lille Grand Palais, May 26-30, 2014.

H. Yoo, R.A. Wibowo, A. Hölzing, R. Lechner, J. Palm, S. Jost, M. Gowtham, F. Sorin,

B. Louis, R. Hock

“Investigation of the solid state reactions by time-resolved X-ray diffraction while

crystallizing kesterite Cu2ZnSnSe4 thin films”

Poster, E-MRS 2012 Spring Meeting, Strasbourg Congress Centre, May 14-18, 2012.

R.A. Wibowo, H. Yoo, A. Hölzing, R. Lechner, S. Jost, J. Palm, M. Gowtham, B. Louis,

R. Hock

“A study of kesterite Cu2ZnSn(Se,S)4 formation from sputtered Cu-Zn-Sn metal

precursors by RTP sulfo-selenization of the metal thin films”

» Best Poster Award «

Appendix

109


R. Lechner, S. Jost, J. Palm, M. Gowtham, F. Sorin, B. Louis, H. Yoo, R.A. Wibowo, R.

Hock

“Cu2ZnSn(S,Se)4 solar cells processed by rapid thermal processing of stacked elemental

layer precursors”

Talk, E-MRS 2012 Spring Meeting, Strasbourg Congress Centre, May 14-18, 2012.

A. Hölzing, R. Schurr, H. Yoo, R.A. Wibowo, R. Lechner, J. Palm, S. Jost, R. Hock

“Real-time investigations on the formation of Cu(In,Ga)(S,Se)2 while annealing Cu-In-

Ga precursors with different sulfur-selenium mixtures”


During the master thesis

Hyesun Yoo, JunHo Kim, Lixin Zhang

“Growth of Inx(S, O, OH)y films by chemical bath deposition”

Poster, International Union of Materials Research Societies - International Conference

on Electronic Materials 2010 (IUMRS-ICEM 2010), KINTEX, Seoul, Korea, August

22-27, 2010.


“Growth of Cu2ZnSnS4 Films by Sputtering with Post-Sulfurization”

Poster, AIP conference, Seoul, Korea, July 25-30, 2010.

Dong-Yeup Lee, Hyesun Yoo, Ki-Bong Song, Jae Ho Yun, JunHo Kim

“Growth of sprayed CIS film and solar cell application”

Poster, 35th

IEEE Photovoltaic Specialists Conference (PVSC), Honolulu, HI, June 20-

25, 2010.


“Comparative Study of Cu2ZnSnS4 Film Growth”

Poster, 19th International Photovoltaic Science and Engineering Conference and

Exhibition (PVSEC-19), ICC Jeju, Korea, November 9-13, 2009.


“Growth of Cu2ZnSnS4 thin films using sulfurization of stacked metallic films”

Poster, the 2nd

International Conference on Microelectronics and Plasma Technology

(ICMAP 2009), BEXCO Convention Center, Busan, Korea, September 22-25, 2009.

Appendix

110

Acknowledgements

First of all, I would like to thank God who had already determined everything before I

come here and start this study. Moreover, without any help or advice of people

who I wrote below, this study would have taken much longer to complete.

Prof. Dr. Rainer Hock who is the Doktorvater of me. Although he has a tremendous

knowledge of crystallography and physics, he is a modest man with a ready wit. I

could learn from him what the scientist is and how one can maintain the balance

between work and one’s life. I am so proud of me that you are my doctor father.

Prof. Dr. Susan Schorr who accepts the review of this dissertation. Thanks for spending

your time to examine my results.

Dr. Robert Lechner who was the supervisor in AVANCIS. Thank you for your support

with full of your kindness for this study: great discussion with you, samples for

metallic layers, Raman scattering data, English grammar for my papers, and so on..

Dr. Gowtham Manohara who was the first supervisor in Saint-Gobain Research (SGR).

Thank you for a glass of beer in Paris together with your advice. That gave me a

confidence for starting this study when I was overwhelmed with this project.

Dr. Arnaud Verger who was the second supervisor in SGR. Thank you for your support

for me and this project and great discussion with you. I also would like to thank to

Virginie Moreau, Corinne Papret, and Francois-Julien Vermersch in SGR for the

great discussion and data for these samples.

Dr. Joao Abreu who was third supervisor in SGR. Although I couldn’t see you during this

study for preparation of several conferences, I would like to thank you that you

hand my papers over to SGR to get a confirmation from them. It might be

wonderful if I could discuss about these data with you, too.

Dr. Stefan Jost in AVANCIS who worked and had great discussions with us for this study.

I would like to thank you along with Dr. Jörg Palm for your support in this study.

It was pity that AVANCIS took a step backward from this project while it was

going on, but the beginning of this project was so wonderful, working with

AVANCIS. Thanks for allowing me to study this subject in Erlangen.

Dr. Dieter Schollmeyer who helped us repair the rotating anode X-ray generator (in-situ

XRD) although he works in University of Mainz. Without your help, it would be

difficult to keep going the measurement for this study.

Dr. Stefan Möckel and Prof. Dr. Peter Wellmann in I-MEET, University of Erlangen-

Nürnberg. Thank you for allowing me to use the evaporator for Se layer on my

samples.

Rameez Ahmad and Thmas Macken who allow me to use Raman spectroscopy when I

urgently needed the scattering data.

I really had a great time staying in LKS (Lehrstuhl für Kristallografie und Strukturphysik).

Although I hesitated to hang out with colleagues at the beginning due to the full

with fear and dread for different languages and cultures, I could feel that the guys

Acknowledgements

111

working in here are so kind and nice. Anyway, consequentially, I met lots of

friends and had wonderful time in here, groups of lunch table, cake, Kartfahren

(go-kart), and so on.. It was really great to meet and know you guys. For you guys,

I could keep going my work and could finish my writing. I would like to thank all

of them who worked in this chair. Additionally, this chair is quite comfortable for

staying overnight. ;-)

Dr. Astrid Hölzing who helped me a lot and cheered me up when I faced with some of

problems and worse situations.

Dr. Rachmat Adhi Wibowo who can speak Korean very well and has lots of knowledge

about kesterite based solar cell. It was really wonderful that I could discuss the

kesterite issues with you in Korean language. -_-b

Marco Brandl who looked over this dissertation and helped me a lot: English grammar,

discussion about kesterite structure, information on the field of solar cell, and so

on..

Sabine Pompetzki who helped me a lot for staying in this chair and for finishing my

writing in Erlangen. It was nice to enjoy lunch with you.

Lisa Lautner who tries to make me hanging out with other people and let me know the

German cultures. It was also great eating lunch together with Ella Schmmidt and

Isabel Schuldes.

Dr. Andreas Schiener who gave me some tips for work and staying in LKS. For your

advices, I did not hesitate to discuss about my results with other colleagues any

more. Additionally, thank you for Kartfahren.

Marvin Beringhof, Patrick Seitz, and Dennis Noll who needed to listen to my unimportant

chat or mumbling because of one reason that they sat next to me. For your listen, I

could have ordinary life in this chair with daily conversation.

Zhen Li who could completely understand the thinking of Asian women in this chair. It

was nice to have a time with you during lunch time and 30th

birthday party.

Dr. Matthias Weisser who introduced me this chair and showed my family the German

culture and life style. Thank you for your kindness for me and advices on the

scientist.

Herbert Lang and Jürgen Grasser (the workshop of this institute) who helped me a lot for

some of important materials made by metals. Without your helps, the

measurements could not keep going on.

Heidrun Brückner who helped me to order some of chemical materials and taught me the

way of dealing with these chemical issues.

Christian Bär who helped me the internet and electrical issues together with electric wire.

Dr. Kaustuv Datta and Haimantee Chatterjee who have lots of wits and cheerful

disposition. The tee which you gave me was so warm to melt my frozen heart. ;)

Torben Schindler and Tilo Schmutzler who accept me as a member of their group.

Dr. Alexander Gröschel, Dr. Christoph Bergmann, and Dr. Michael Klimczak who invited

me in their party.

Appendix

112

Without normal life, without good results in one’s work. Therefore, I would like to thank

to people who have no connection with this study: my family and friends. Without

them, I would have nostalgia for Korea so often. Therefore I also appreciate to

Viber and Kakaotalk which may contact with my family and Korean friends.

Dr, Chenyong Si, Dr. Nooshin mir and her husband Dr. Mosoud Azadi, Yan Zhuang,

Valentina Miguez Pacheco, and Dr. Modhaffar Husni Ali who I met in Fürth. It

was really great to meet you guys when I came here in Germany and wonderful to

spend a time with you guys. I would never forget the food which is made by

Chenyong. For you guys, I could feel that I am not alone in Germany at the

beginning of my work.

Dr. Junghyun Lee and Yoonkyung Nam who have lots of chat with me when I have some

problems about my life or my work. It was great time to spend a time with you

guys.

Die Koreanisch Katholisch Gemeinde which has mass one time per month in Korean. The

people who I met in this Gemeinde were really nice and kind. They were really

good persons to share our life and chat each other. Furthermore, the Korean food

which I ate in there was really delicious as much as I didn’t miss Korean food at

all.

The group of Friday dinner table and the group of Korean scientists in Erlangen. It was

really wonderful to have a chat and discuss about lots of issues, not only Korea but

also history and so on.. Talking in mother language makes person release from

shrunk heart and makes person feel comfortable.

The friends who live in Korea. Although I could not contact with you so often as much as

I was in Korea, you always gave me an emotional support. Thank you guys and

see you in Korea again!

Lastly, I would like to thank my parents, Soon Duk Kim (mother) and Sung Choon Ryu

(father), and my little brother, Bumseok Yoo, who I may stay in Germany without

any anxiety about my family. I love you.

When I came here, I had no friend in Erlangen. However, now I have lots of friends and

had great times with them. It was wonderful experiences for me, and I would never

forget it and you guys.

Though thy beginning was small, yet thy latter end should greatly increase.

reaction mechanism between cu zn sn se components for …mechanism... · 2.2 ellingham diagram for...

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