The single phase region in Cu2ZnSnS4 thin films from theory and

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Cite This: Chem. Mater. 2018, 30, 4624−4638

The Single Phase Region in Cu2ZnSnS4 Thin Films from Theory and Combinatorial Experiments Alexandra Davydova, Katharina Rudisch, and Jonathan J. S. Scragg* Ångström Solar Center, Solid State Electronics, Uppsala University, Box 534, SE-751 21 Uppsala, Sweden

Chem. Mater. 2018.30:4624-4638. Downloaded from pubs.acs.org by KAOHSIUNG MEDICAL UNIV on 10/05/18. For personal use only.

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ABSTRACT: Cu2ZnSnS4 (CZTS) is hoped to be a future, earth-abundant absorber material for thin film solar cells, but performance remains below the level needed for commercialization. In this work, the size of the single phase region of CZTS obtained from thin film synthesis methods is explored, to determine the scope available for defect engineering and thereby future improvements in solar cell performance. A chemical model for the single phase region is presented, based on equilibria between defect complexes in the CZTS phase and the other solid- and gas-phase components present during synthesis. The model predicts a variable single phase region size, depending on the partial pressures of SnS and S2. The model is verified by analysis of combinatorial thin-film CZTS samples prepared with different synthetic conditions and characterized by Raman and compositional mapping. We conclude that typical synthesis strategies for CZTS are not capable of accessing the full range of the CZTS single phase region since the required partial pressure of S2 is very large. The important implication is that our understanding of CZTS defect chemistry from experimental studies is incomplete and that scope exists for tuning the defect properties toward better solar cell performance.

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electrostatic potential fluctuations,5 as well as affecting the formation probability of other point defects.6 Defects can also have indirect effects on optoelectronic properties, for example, by causing a change in bulk diffusivity that in turn influences kinetically limited processes such as cation ordering and grain growth.7,8 In recent years, the efficiency of kesterite-based solar cells has reached a bottleneck that appears to relate to poor bulk properties in the absorber layer. The actual causes are still under debate, but the main candidates are deep defects and/or band tailing due to a high density of shallower defects or defect complexes.9,10 It is to be hoped that a better understanding and control of defects in kesterite (“defect engineering”) will help to reach higher solar cell efficiencies or to better understand the intrinsic limitations of the material. The most fundamental way to conduct defect engineering in such a material is by careful control of its composition. In this work, we focus mainly on the sulfide compound CZTS, but the approach and conclusions are transferable to the selenide and mixed sulfo-selenide compounds. For reference, Figure 1 shows two ways to represent the composition space of CZTS. The pseudoternary phase diagram in Figure 1a describes the composition assuming that S content is fixed by the charges on the metal cations. Phase boundaries are deliberately not shown, but an example for the case of equilibrium at 400 °C can

INTRODUCTION During recent years it has become increasingly clear that photovoltaic solar energy technologies (“solar cells”) will become an integral part of global electricity generation, and production of such technologies is anticipated to undergo continued and dramatic growth.1 For this reason, there is great interest in new solar cell materials that can reduce the energy input in solar cell manufacture and that consist of only abundant and cheap components. One of these candidates is Cu2ZnSn(SxSe1−x)4, known either as “kesterite”, for its crystal structure, or abbreviated to CZTS (for x = 1), CZTSe (for x = 0), or CZTSSe (0 < x < 1). The kesterites are nontoxic materials with band gaps from 1.0 eV (pure selenide) to 1.5 eV (pure sulfide)2 that make them suitable for absorber layers in thin film solar cells. The best reported kesterite devices have power conversion efficiencies in the range 11−13%.3,4 However, to reach the goal of >20% efficiency needed for realistic commercial technology, a better understanding of the limiting factors is needed. One issue in particular that is poorly understood is the relationship between the composition of the kesterite phase and the resulting optoelectronic properties of the film. The fundamental relationship between composition and properties is that the former affects the occurrence and concentrations of point defects and donor−acceptor defect complexes. Intrinsic point defects are to a large extent responsible for the doping and recombination characteristics in kesterites, and intrinsic defect complexes of different kinds, present in large concentrations, can lead to net band gap variations and localized band gap or © 2018 American Chemical Society

Received: March 22, 2018 Revised: June 20, 2018 Published: June 20, 2018 4624

DOI: 10.1021/acs.chemmater.8b01213 Chem. Mater. 2018, 30, 4624−4638

Article

Chemistry of Materials

for the observed “optimal” composition for kesterite absorber layers can be found from first-principles defect calculations, with Cu-poor, Zn-rich compositions thought to lead to beneficial doping levels and to reduce the probability of forming detrimental intrinsic defects such as SnZn and the complex {2CuZn− + SnZn2+}0. In terms of the defect “types” A−J as illustrated in Figure 1 (see also Table 1), it is believed on the basis of the apparent optimal composition that A type, where the defect complex {VCu− + ZnCu+}0 occurs, is most favorable for solar cells.14 Table 1. CZTS “Types” and Associated Neutral Defect Complexes Detected or Proposed in CZTS and/or CZTSe To-Date, According to Refs 15,24−27a type

defect complex

category

A D B C E F G H I J (K) (L)

{VCu− + ZnCu+}0 {Cui+ + CuZn−}0 {2ZnCu+ + ZnSn2−}0 {2CuZn− + SnZn2+}0 {2VCu− + SnZn2+}0 {2Cui+ + ZnSn2−}0 {ZnSn2− + Zni2+}0 {SnZn2+ + VZn2−}0 {CuSn3− + 3Cui+}0 {SnCu3+ + 3VCu−}0 {CuZn− + CuSn3− + 4Cui+}0 = I + D {ZnCu+ + SnCu3+ + 4VCu−}0 = J + A

{v} {i} {s} {s} {v} {i} {i} {v} {i} {v}

a

The categories in the third column are assigned to complexes containing vacancies ({v}), interstitials ({i}), or only substitutional defects ({s}), respectively.

Unfortunately, the observations from experiments do not necessarily relate to fundamental properties of CZTS. This is because, in the vast majority of cases, it is the integral film composition that is reported and discussed, rather than the CZTS phase composition. The integral composition includes contributions from both the CZTS phase and the multiple secondary phases that can be present alongside, detected or otherwise.15 It is widely acknowledged that single phase CZTS is hard to produce, and reliably determining the actual CZTS phase composition in this situation is not trivial. This is problematic from a defect engineering perspective since it is the CZTS phase composition that is important in determining the defect makeup. As a result, while certain basic effects of composition can be rationalized, we still do not have a clear and accurate picture of the relationship between CZTS composition and the important material properties that influence device performance. To get further, we need to first gain a better idea of the limits of off-stoichiometric compositions that can be reached for the CZTS phase itself in thin films. What this amounts to is determination of the extent of the single phase region or homogeneity range of the CZTS phase that can be obtained in thin film synthesis. Given better information about this, we can ascertain not only the range of compositions for which secondary phases are avoided, we can also determine the “process window” for defect engineering since the extent of the single phase region determines the total possible scope for varying the composition of the CZTS phase and thereby changing its defect characteristics. In this study, we use combinatorial thin films to investigate the single phase region of CZTS. Combinatorial methods can be

Figure 1. Two ways of representing the composition of a CZTS sample: (a) a pseudoternary phase diagram and (b) a metal-ratio diagram. The lines A−J represent CZTS containing different types of defect complexes (see text for further explanation).

be found in ref 11 (or ref 12 for CZTSe). The interesting range of compositions for solar cells is at most 10 at. % from the stoichiometric point, and therefore, it is often more convenient to magnify this area. When the S content is undefined and for ease of plotting, a figure such as Figure 1b is often used, where composition is represented via two ratios of the metal atom concentrations. In this Article, we use the ratios Cu/Sn and Zn/ (Cu + Sn), which take values of 2 and 1/3 for perfect stoichiometry. For simplicity, we use the terms “Cu-rich” and “Sn-rich” to refer to compositions with Cu/Sn > 2 and Cu/Sn < 2, respectively, and the terms “Zn-rich” and “Zn-poor” to refer to compositions with Zn/(Cu + Sn) > 1/3 and 0.33. The poor performance of Cu-rich material is thought to relate to the presence of Cu−S secondary phases with narrow band gap and semimetallic properties. On the Sn-rich side, the secondary phases seem less problematic.13 Further justification 4625


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stoichiometry and device performance. The films were prepared by cosputtering of CuS, SnS, and ZnS targets followed by sulfurization at 550 °C. In this process, Cu/Sn ratios on both sides of stoichiometry were obtained (both Sn-rich and Curich), but Zn-poor compositions were not captured, so the stoichiometric point was not within the sample area. Solar cell structures were made and the measured solar cell performance varied strongly across the compositional range. In contrast to the findings of Fairbrother et al., the best overall power conversion efficiency (7.5%) was now obtained along the Cu/Sn = 2 stoichiometric line, however again there were separate optima for Voc and Jsc, the former toward the Zn-rich side, and the latter in the region with Cu/Sn < 2. Collord et al.23 made an interesting and detailed study of the optoelectronic properties of combinatorial sulfo-selenide CZTSSe, produced by annealing films from nanoparticle inkbased precursors. Their results showed a dramatic difference in the quasi-Fermi level splitting as Cu-content varied, showing superior optoelectronic quality for material formed with a Cupoor, Zn-rich integral composition. The Sn content was not found to be important, but similarly to the work of Du et al., the films experienced Sn-loss during annealing. Since secondary phases were not characterized in this study, the actual variation of Sn content in the CZTS phase, and therefore any role in the optoelectronic behavior, could not be fully clarified. These reports exploited the combinatorial method to highlight the strong influence of composition on CZTS properties, including resultant device performance. However, as noted the composition of the CZTS phase at each point on the sample is masked by the presence of secondary phases, and no single phase regions were explicitly identified. The work so far also shows that there is interplay between the synthesis conditions chosen and the composition range of the CZTS phase that can be obtained (most striking in the case of Du et al.), which suggests that the partitioning of material between the CZTS phase and secondary phases, i.e., the size of the single phase region, might vary depending on the preparation procedure. Thus, before answering the question of which composition is preferable for CZTS solar cells (and, following that, why), we should first determine which CZTS compositions are possible under given synthesis conditions, and, as an extension, which compositions might be possible by tuning the synthesis in new ways. As we will show here, the single phase region is not fixed, even for a fixed temperature, but depends on the conditions of synthesis. We propose a theoretical framework to describe the single phase region based on equilibria of defect complexes, and go on to validate this experimentally by investigating CZTS properties and secondary phase distributions in combinatorial CZTS thin films made using different conditions.

efficient for studying large compositional ranges while keeping all other conditions constant. The type of combinatorial method used in thin film contexts is usually a “continuous composition spread” or CCS method,16 whereby thin films are deposited with smooth gradients in composition over their surface in one or more directions. This is typically done in a cosputtering or coevaporation system using two or three material sources focused onto a nonrotating substrate. Samples are then analyzed by mapping or line-scanning with various techniques, the results of which are correlated to the composition at each probed location on the sample. Several research groups have used such combinatorial approaches to gain information about the relations between the composition of the CZTS film and resulting properties. Du and Lund et al.17,18 published two papers concerning structural and photoluminescence properties of CZTS as a function of stoichiometry using combinatorial CZTS films prepared by coevaporation at different substrate temperatures. A wide composition range could be captured on a sample deposited at a temperature of 325 °C. For films synthesized at 470 °C, the highest temperature used in that study, the compositions on the side of the sample that received greatest Sn flux during deposition nevertheless collapsed to the Cu/Sn = 2 (B−C type) line, i.e., Sn-rich CZTS could not be formed. This was attributed to the volatility of Sn−S phases and Sn-rich ternary/ quaternary phases during the high temperature growth under vacuum, a well-known issue for CZTS for such “single-stage” synthesis methods.19 The authors were able to observe many secondary phases using Raman scattering, although the single phase region was not defined due to reasonable doubts about the detection limits of this technique. The CZTS on the side of the sample that received the highest Sn flux during growth showed low photoluminescence yield and lower structural quality, based on Raman peak width. Although the material in this region was not actually Sn-rich, these observations were attributed to the presence of Sn-related point defects. Meanwhile, the Cu-rich side of the sample showed much more intense luminescence and sharper Raman peaks. A feature of note on the samples was a specular line coinciding with the Cu2SnS3−CZTS−ZnS tie-line (where the ratio Cu/Sn = 2), which is also the B−C line shown in Figure 1. Fairbrother et al. prepared a combinatorial thin film of the selenide CZTSe compound, using cosputtering of metallic precursors followed by selenization at 550 °C.20 The resulting sample contained compositions on all sides of the stoichiometric point. After processing the film into nearly 200 individual solar cells, maps of photovoltaic parameters as a function of composition were produced. The Cu-rich (Cu/Sn > 2) devices were all poor, but good efficiencies were obtained on the Sn-rich side (Cu/Sn < 2), opposite to the photoluminescence behavior observed by Du et al., but explainable by the influence of Cu−S secondary phases on the Cu-rich side. Interestingly, it was found that within the Sn-rich range, there were two different “optimal” compositions, where either the open circuit voltage (Voc) or the short circuit current density (Jsc) were maximized. The best overall efficiency of 6.9% was found for Cu-poor, Zn-rich integral compositions near the “A-type” line. The same sample was studied in detail using multiwavelength Raman scattering, and relationships were determined between certain spectral features and the CZTS composition, caused by the presence of neutral defect complexes of various kinds.21 Valentini et al.22 also used combinatorial thin films to investigate the correlation between CZTS (sulfide) film

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THEORY OF THE SINGLE PHASE REGION The single phase region (SPR) of CZTS is the range of compositions for which the crystal structure (kesterite or disordered kesterite) is stable with respect to precipitation of secondary phases, for a given set of conditions. Somewhere within the SPR is the stoichiometric point, where perfect kesterite with composition exactly represented as Cu2ZnSnS(e)4 is obtained. At all other compositions within the SPR, the kesterite phase forms with a collection of defects that make it nonstoichiometric. In the equilibrium phase diagrams,11,12 the SPR extends toward Zn-rich/poor, Cu-rich, and Sn-rich 4626


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This makes complex formation energetically favorable overall, but only up to a point. At a certain concentration of complexes, the entropy gain ceases to outweigh the enthalpic penalty of adding them to the lattice. Above this concentration, increasing the availability of Zn(S), in this example, will only result in formation of ZnS secondary phase particles. At a given temperature, the concentration limit or “solubility limit” of Btype complexes in CZTS (which could be expressed as a concentration of complexes per unit volume, for example) defines one point on the SPR boundary, specifically the point where the SPR intersects the B-type line on the phase diagram. The extent of the SPR along the other lines is determined by the solubilities of each of the defect complexes concerned (see Table 1). The solubility limits will depend on synthesis conditions, especially temperature. For this reason, the single phase region must be determined for the temperature of interest, which is around 450−550 °C for CZTS(e) thin films. Most information about the SPR to-date comes from solid-state synthesis at much higher temperatures11,12,24−29 and thus cannot be directly applied to the thin film case. Besides temperature, it is important whether the synthesis is at equilibrium or not. The equilibrium case, used for construction of phase diagrams, gives certain values of the solubility limits, but these values do not necessarily apply to thin film synthesis. In an equilibrium synthesis, the temperature is well-defined, and the chemical potentials, determined by the proportions of the solid components (e.g., Cu, Sn) and the partial pressures of gas phase components (S, Se, Zn, SnS etc.), are fixed. High temperatures and extended time periods are used to ensure equilibrium is reached. After that, rapid cooling is used to freeze the equilibrium phase mixture in place for subsequent characterization. With few exceptions, solar-cell relevant material is synthesized under nonequilibrium conditions. The partial pressures of gas phase species, i.e., their chemical potentials, may change during synthesis and may be much lower or higher than the equilibrium case. The synthesis temperature is likely lower and the allowed time scales much shorter, so that equilibrium may not actually be reached even if the chemical potentials were fixed. Finally, the cooling rate may be relatively slow, allowing reactions to occur in response to changing conditions during the cool-down phase. None of this means that we cannot analyze material formation in nonequilibrium synthesis; it only means that the equilibrium SPR, even if it were available for the synthesis temperature of interest, does not provide a complete picture. To understand how the solubility limits of defect complexes, and thereby the SPR size and shape, could be affected by synthesis conditions, we first need to derive formation reactions for the defect complexes. We do this by considering an equilibrium in which additional elements are partitioned between the kesterite phase and secondary phases. For example, at an interface between a kesterite grain and a ZnS (sphalerite) grain, it may be energetically favorable for some Zn and S atoms to cross the interface and join the kesterite phase. This could be described as “dissolving” of ZnS into the kesterite lattice to form Zn-rich CZTS (B-type). If the conditions are changed (e.g., a lower temperature), then the reverse process might become favorable; ZnS would instead “precipitate” from the kesterite phase, which would have a lower Zn(S) content as a result. In fact, all of the defect complex types in Table 1 are compositionally equivalent to dissolved secondary phases in appropriate amounts; this can be easily seen for B/C/D/E-type complexes since their “lines” coincide with the CZTS-ZnS/ Cu2SnS3/Cu4SnS4/SnS2 tie lines; see Figure 1a. The other

compositions (i.e., to all sides of the stoichiometric point in Figure 1. For a material to have an extended SPR, it must be possible to introduce defects with sufficiently low formation energy. In CZTS(e), there is an abundance of neutral defect complexes (pairs or triplets of point defects with compensating charges) that enable nonstoichiometry. These have been given labels “A”, “B”, ..., “L” in the order in which they were first determined. The neutral defect complexes that have been proposed and, in most cases, experimentally identified in CZTS(e) are listed in Table 1. Addition of any of the complexes in Table 1 to the CZTS(e) lattice will change the stoichiometry without affecting net doping. In Table 1, the defects are grouped in pairs with opposing compositional influences, e.g., B-type gives pure Znrich composition, whereas C-type gives pure Zn-poor composition; note also that the component point defects within each pair are “opposites”, e.g., Cui vs VCu, SnZn vs ZnSn, etc. The proposed K and L type complexes are in parentheses because these complexes are simply equivalent to a combination of two other complex types, (K = I + D and L = J + A) and therefore do not represent unique, “irreducible” types. Thus, we do not consider them further here. We note that in this Article we only consider neutral defect complexes and not isolated point defects. Complexes, by definition, have a lower formation energy compared to the isolated point defects and are therefore present in higher concentrations. Thus, while isolated point defects are critical for the opto-electronic properties of CZTS, in this work we make the assumption that it is the defect complexes that play the dominant role in determining the stoichiometry of the CZTS phase and the size and shape of the SPR. This work constitutes a first step in which we define how the limits of the SPR depend on synthesis conditions. In the future, we aim to study how the opto-electronic properties of the CZTS phase are influenced by the point defects formed against the “background” of more abundant defect complexes across the single phase region. Since there are so many of them, it is convenient to categorize the defect complex types. We do this according to whether they contain vacancies, interstitials, or only substitutional defects. Thus, the vacancy-containing complex types A, E, H, and J are collectively referred to as {v} complexes, the interstitialcontaining complexes D, F, G, and I are referred to as {i} complexes, and the remaining B and C types, which contain only substitutional defects, are referred to as {s} complexes (see Table 1). Figure 1b indicates the regimes of CZTS phase composition where these different categories would be found (if the given composition was part of the SPR). One can see that {v} complexes are present in the range where Cu/Sn in the CZTS phase is less than 2 and {i} complexes are present when Cu/Sn in the CZTS phase is greater than 2. When Cu/Sn is exactly equal to 2, only {s} complexes are present. Then there are regions in which two categories are present, labeled {v} + {s} and {i} + {s}. The concentrations of the defect complexes that arise in a given sample, and therefore the stoichiometry of the CZTS phase, depend on the availabilities (formally, the chemical potentials) of the different elements during synthesis, as well as the synthesis temperature. For example, if excess Zn(S) is available during synthesis, then some of this may be incorporated in the CZTS phase, via the formation of B-type defect complexes. Although the formation enthalpies of the defect complexes are positive,5 their presence increases the entropy of the crystal, increasingly so at higher temperatures. 4627


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original empty CZTS unit may remain, also with an effective charge. Note that the vacancies do not need to exist “from the start” within the CZTS lattice, but they can be created at a surface or interface during the reaction. It is assumed that the resulting defects, which are by necessity of equal and opposite charge, form complexes due to their mutual electrostatic attraction. Indeed, this complexation is actually required to reach an appreciable defect concentration (and thereby nonzero SPR) since many of the point defects in isolation (e.g., Cui, VZn, etc.) have very high formation energies.5 Following this line of reasoning, we can write straightforward formation reactions for each of the defect complexes, with atoms moving between a secondary phase and the CZTS phase. In the below reactions, the notation [A], [B], etc., indicates a formula unit of CZTS containing the given defect complex. In the forward direction, the reactions correspond to dissolving of secondary phases in CZTS, and in the reverse direction, they correspond to precipitation of secondary phases.

defect complexes require a mixture of two secondary phases to be considered. This equivalence arises simply because the point defects have the same charges as the cations in the stoichiometric phases, and the defect complexes must be neutral. It is possible to imagine other neutral defects such as Cu(II)Zn and Sn(II)Zn. These have not been investigated in any detail so we exclude them from our model; however, in the Results and Discussion section, we elaborate further on the possible presence of Cu(II)Zn. In the following, we describe a way to visualize the equilibria between defect complexes and secondary phases; the same ideas are illustrated schematically in Figure 2. First, we consider an

2ZnS(s) + SnS2(s) V [A]

(1)

Cu4SnS4 (s) V [D]

(2)

4ZnS(s) V [B]

(3)

4Cu 2SnS3(s) V [C]

(4)

2SnS2(s) V [E]

(5)

2Cu 2S(s) + 2ZnS(s) V [F]

(6)

Cu 2S(s) + 3ZnS(s) V [G]

(7)

2Cu 2SnS3(s) + SnS2(s) V [H]

(8)

3Cu 2S(s) + ZnS(s) V [I]

(9)

2ZnS(s) + 3SnS2(s) V [J]

(10)

From this basis, we can consider how changing synthesis conditions might influence the SPR by altering the state of the equilibria above. Since all these reactions contain only solid phases, the equilibrium constants that determine the partitioning of atoms between defect complexes and secondary phases will only depend on temperature. This would mean that the single phase region had a single, defined shape and size for a given synthesis temperature. However, the reactions above do not fully capture the situation in (thin film) synthesis of CZTS. What they omit is that SnS2 is not stable in the solid state for a wide range of the synthesis conditions often used. Indeed, SnS2 is rarely observed in thin film synthesis because it decomposes as below:19

Figure 2. Schematic illustrating formation of defect complexes by “dissolving” various secondary phases in the CZTS lattice. The vacancies in an “empty” CZTS formula unit are filled with the atoms from the given secondary phase(s), resulting in multiple defects with net zero charge. For H/J- and C-type, two and three formula units of CZTS are required, respectively.

empty formula unit of CZTS, i.e., the vacancies {2VCu−, VZn2−, VSn4−, 4VS2+}. Then, we take enough formula units of a secondary phase to provide four sulfur atoms (i.e., 2SnS2 or 4ZnS, etc.). Then, to the greatest extent possible, we fill the vacancies “correctly”, using the atoms of the secondary phase(s), thus forming, e.g., SS, CuCu, ZnZn, etc. In doing so, we neutralize the effective charges of the respective vacancies. The remaining atoms from the secondary phase(s) can either occupy substitutional positions (e.g., ZnSn2−) or go to interstitial sites (e.g., Cui+), whichever has the lowest formation energy in the given case. In these positions, the atoms are now defects and have an effective charge. Finally, unfilled vacancies from our

SnS2 (s) → SnS(g/s) +

1 S2 (g) 2

(11)

SnS can be in the gas phase alone or in both solid and gas phases, depending on its partial pressure. For conditions where SnS2 is not stable, the defect complex formation reactions that involve SnS2 are no longer valid and should be rewritten in terms of the gas phases: 2ZnS(s) + SnS(g/s) +

1 S2 (g) V [A] 2

2SnS(g/s) + S2(g) V [E] 4628

(12) (13)


Article

Chemistry of Materials 2Cu 2SnS3(s) + SnS(g/s) +

1 S2 (g) V [H] 2

2ZnS(s) + 3SnS(g/s) + 1.5S2(g) V [J]

experimental results that validate the hypotheses described above. So far we have not considered anion-related defect complexes, containing, e.g., S vacancies (VS), for the simple reasons that these have not been experimentally detected. While one might expect that VS would become important under S-poor conditions, e.g., annealing with low or zero S partial pressure, the above equilibria show that actually it is possible for CZTS to lose S without the formation of VS. As we will see, the above reactions seem sufficient to describe observed behavior in our experimental samples. If VS point defects were to be experimentally detected, possibly in complexation with Sn(II)Zn,30 the model above could be extended appropriately.

(14) (15)

In Reactions 16−18 we include the additional possibility that CZTS types can be interconverted via exchange of SnS and S with the gas phase, specifically that the Zn-poor C- and H-types can be interconverted and that the Zn-rich B-, A-, and J-types can be interconverted, depending on the amount of SnS and S available: [C] + 2SnS(g/s) + S2(g) V 2[H]

(16)

[B] + 2SnS(g/s) + S2(g) V 2[A]

(17)

[B] + 6SnS(g/s) + 3S2(g) V 3[J]

(18)

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EXPERIMENTAL SECTION

In order to explore a broad range of CZTS compositions while keeping all other conditions constant, compositionally graded thin films were prepared. The preparation process is in most respects the same as our baseline CZTS solar cell production process at Uppsala University.31 First, fully sulfurized but poorly crystalline “precursor” Cu−Zn−Sn−S films are produced by reactive cosputtering, and in a second step, these are recrystallized in a tube furnace. Sputtering was done in a Von Ardenne CS 600 deposition chamber. The base pressure of the system is below 1 × 10−6 mbar. Targets of CuS, Zn, and Sn were cosputtered in a H2S and Ar atmosphere supplied at mass flow rates of 10 and 30 sccm, respectively, providing constant total pressure of 6.65 × 10−3 mbar. The target powers were set to 130, 440, and 90 W with 20 kHz pulsing frequency for Zn and Sn and 150 kHz for CuS. Substrate heating was set at 375 °C yielding a temperature of the substrate of about 225 °C. The sputtering time was set to 25 min giving a film of approximately 500 nm in thickness. For compositionally graded samples, substrate rotation is switched off, whereupon compositional gradients arise naturally across the area of the sample (initially 70 × 70 mm), reflecting the geometry of the sputtering targets and substrate holder and the variations in sputter flux received by each point on the sample. The targets are distributed at three corners of a triangle pointing at the center of the substrate holder, at approximately 160 mm distance. In this way, each precursor contains compositions spanning a large enough range to capture the single phase region for CZTS. Initial precursors were deposited on Mo-coated sodalime glass slides of 70 × 70 mm size. To determine the extremes of the composition range, 10 × 10 mm pieces from each corner of the samples were removed and measured with XRF calibrated by RBS. Based on these measurements, the remaining samples were cut to 50 × 50 mm size for further processing. Two such samples, C1 and C2, are discussed in this Article. The recrystallization process was carried out in a tube furnace at 550 °C under a static argon atmosphere (350 mbar). Each sample was loaded into a graphite box together with additional S (and Sn) sources (see below). The samples were then heated up from room temperature up to 550 °C with a rate of 4 °C s−1, by insertion of the box into a preheated hot zone. After 10 min, the box was extracted to a cold zone and cooled down to 200 °C within 2 min and from 200 °C to room temperature with an average cooling rate of about 4 C min−1. For sample C1, the precursor was loaded together with 75 mg of elemental sulfur (99.999%, Alfa Aesar). This corresponds to our baseline solar cell process, for which SnS secondary phases, but not SnS2, are observed.13 For sample C2, the box was prepared with an in situ synthesis of SnS2 prior to precursor recrystallization. To achieve this, approximately 40 mg of tin(II) chloride (99.99%, Sigma-Aldrich) and 40 mg of elemental sulfur were loaded in the graphite box, which was then taken through the above-described heating process (without the precursor). SnCl2 has an extremely high vapor pressure, around 400 mbar at 550 °C,32 and along with the provided S, bright orange deposits of pure SnS2 are formed in the box, confirmed using Raman spectroscopy. Subsequently, the precursor C2 was loaded, along with additional S (75 mg) and the recrystallization process was carried out. In this process, the large amount of SnS2 and residual SnCl2 keep the atmosphere saturated with both Sn-bearing molecules (SnS and SnCl2) and S, resulting in SnS2

The above equilibria Reactions 12−18 only affect the {v} complexes, those containing vacancies, which occur on the Snrich side of the phase diagram (Cu/Sn < 2). They show that the concentration of {v} complexes in the CZTS phase at equilibrium will depend not only upon the temperature but also upon the partial pressures of SnS and S2, and specifically on the product PSnSPS21/2. The lower the value of this product, the lower the solubility of the relevant defect complexes and the more secondary phases will appear for a given integral composition. Based on these results, we can make the following testable hypotheses: (1) The maximum extent of the SPR toward Cu/Sn < 2, i.e., the maximal concentrations of {v} complexes A, E, H or J − will be observed under conditions in which SnS2 is a stable solid phase (high S2 and SnS partial pressures). This is the case in the equilibrium phase diagram of Olekseyuk et al. determined from solid-state synthesis11 but is usually not so in thin-film synthesis. (2) The extent of the SPR toward Cu/Sn < 2 will decrease for lower partial pressures of S2 and SnS; in other words, it will be harder to incorporate {v} complexes even if the integral composition is Sn-rich (for example, in annealing of a Sn-rich Cu−Zn−Sn precursor); secondary phases of SnS will instead precipitate and potentially evaporate. (3) In the absence of either S or SnS partial pressures, the SPR will have zero extent toward Sn-rich compositions: {v} complexes will not be formed, and all Sn excess will be contained in secondary phases regardless of the integral composition (as long as the integral composition does not by itself introduce sufficient partial pressures of S and SnS into the gas phase). (4) In contrast, the solubility limits of {s} (B- or C-type) and {i} (D-, F-, G-, or I-type) complexes, and the resulting SPR limits toward Zn-rich/poor and Cu-rich compositions should not depend on the partial pressures in the gas phase, since the formation reactions of the relevant complexes only contain stable solid phases. As far as defect engineering is concerned, therefore, we should be aware that the extent of the SPR, and thereby the accessible range of material properties in vacancy rich CZTS, i.e., the range usually targeted for solar cells, will depend very much on the conditions provided during material synthesis. The question arises as to what extent we are able to produce the desired CZTS compositions, and whether we have complete access to the composition range in which properties are most suited to solar cell functionality. In the following sections, we present 4629


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Chemistry of Materials being kept stable in the solid phase for most of the anneal. We have also tried adding SnCl2 and S directly together with the precursor, but in this case, the sample becomes covered with SnS2 precipitates, presumably from a gas phase reaction between SnCl2 and S, which interfere with subsequent analysis. After annealing, mechanical scribing was used to define an area on each sample of approximately 30 × 40 mm. The scribe lines were used to align the various mapping measurements described below so that all measured data could be treated in a consistent coordinate system. To enhance the Cu−Zn ordering in the resulting samples, they were in each case subjected to additional low temperature treatments consisting of a fast ramp to 250 °C and a slow cool to room temperature at a rate of 0.1 °C min−1 (after33). The composition ranges on the graded films were measured by energy dispersive X-ray spectroscopy (EDS) mapping with a Zeiss Leo 1550 system, using an accelerating voltage of 20 keV. Area measurements were made on a grid within the scribed border, with each “pixel” having dimensions of approximately 2 × 1.5 mm. Only the cation contributions were taken into account due to overlap of S−K with Mo−L lines. In quantification, Cu−K/L, Zn−K/L, and Sn−L lines were analyzed. The chosen accelerating voltage ensures that the full depth of the film is analyzed. We emphasize that the EDS measurement gives the integral composition (CZTS + secondary phases) and not the composition of the CZTS phase itself, and it is further assumed that there are no composition gradients within the film depth. The ability of EDS to provide accurate composition values can be questioned, but it is hard to cross-check the combinatorial samples due to the composition gradients. From experience, we find that the Cu/Sn ratio given in our EDS system is generally comparable to that measured using calibrated XRF measurements but that the Zn content in the sample is often slightly overestimated by EDS. In this study, we focus less on the absolute values of composition and more on the trends observed within and between the two samples; in this case EDS measurements are sufficient. The overall integral composition ranges for the samples C1 and C2 are shown in Table 2. In both cases, the stoichiometric point is

phase itself. Finally, we discuss the observations in the light of the theory introduced above. Part I: Analysis of Secondary Phase Distribution. Raman mapping data for excitation wavelengths of 325, 532, and 785 nm were analyzed to determine the distributions of certain secondary phases on the samples. In the mapping experiments, the observed phases besides CZTS were Cu3SnS4 (532, 785 nm) and ZnS (325 nm). The other Cu−Sn−S ternary phases, Cu2SnS3 and Cu4SnS4, could not be resolved in any of the measurements. This may be due to weak Raman signals/ overlap with other phases (notably CZTS itself), or it may just be the case that the Cu3SnS4 was the most stable ternary under the given synthetic conditions and the others did not form. Sn−S phases were only detected by more detailed point analysis with Raman, and Cu−S phases by SEM-EDS point analysis. Here, we present processed Raman data in the form of intensity maps. In the Supporting Information, examples of individual spectra are given. The maps in Figure 3 show the UV−Raman intensity integrated in the range 680−730 cm−1 for samples C1 and C2, as a function of integral composition. This spectral range captures the second-order mode for ZnS and is mostly free of interference

Table 2. Integral Composition Ranges Captured in the Two Combinatorial Samples C1 and C2 sample C1 C2 stoichiometry

Cu/Sn 1.58 1.55 2.00

Zn/(Cu + Sn) 2.40 2.45

0.30 0.28 0.33

0.54 0.49

located within the sample but toward one side, so that the range of Znpoor compositions analyzed is rather small. However, secondary phases could be observed on all sides of the stoichiometric point (see later), indicating that the SPR itself should be fully captured. Multiwavelength Raman spectroscopy was performed in a Renishaw inVia system with excitation wavelengths 325 nm (for analysis of ZnS phases at the surface), 532 nm (for analysis of Sn−S, Cu−S, Cu−Sn−S, and CZTS phases), and 785 nm (for resonance-enhanced analysis of the CZTS phase). The grid dimension for mapping was 2 × 2 mm within the same area as for EDS. For each sample C1, C2 composition mapping and multiwavelength Raman mapping were performed both before and after ordering procedure, from which we could conclude that the ordering procedure does not influence composition or the distribution of secondary phases. The collected data were analyzed and combined into compositional maps using an in-house R code.

■

RESULTS AND DISCUSSION This section is split into three parts. First, we describe the distributions of secondary phases on the samples, which allow us to identify the “outer limits” of the single phase region (SPR) in the respective samples. Second, we analyze the resonance Raman of the CZTS phase, which provides a way to see how the properties of the CZTS itself vary with composition and thereby obtain information on the compositional variations in the CZTS

Figure 3. Maps of the 325 nm Raman intensity in the range of the ZnS second order mode (680−730 cm−1) as a function of integral composition, for samples C1 (a) and C2 (b). The dotted line corresponds to the boundary between the presence and absence of observable ZnS peaks in the spectra, from visual inspection. The intensity scale is in arbitrary units, and values cannot be compared from C1 to C2. 4630


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Chemistry of Materials with the CZTS peaks.34 Judgement of the presence/absence of ZnS was made by visual inspection of the spectra. The dashed lines in the figures show the boundaries between the regions where ZnS is observed (above the line) and not observed. As expected, ZnS is only seen on the Zn-rich sides of the stoichiometric point. In both cases, the boundary of observable ZnS is in the range Zn/(Cu + Sn) = 0.36−0.4, which allows for a nonzero single phase region of CZTS toward Zn-rich stoichiometry in both samples. In the Raman measurements made using the 785 nm laser, CZTS and Cu3SnS4 (CTS) are both in resonant conditions. The main CTS mode is at about 318 cm−1,35 which means that there is overlap with the “tails” from the CZTS main mode at 338 cm−1 and the mode at 305 cm−1.36 Figure 4 shows maps

absence of CTS, and the associated boundaries are shown by the dotted lines in Figure 4 (CTS below the lines). CTS is tightly confined to the Zn-poor corner of the samples, and the boundaries nearly touch the stoichiometric point, indicating that substantially Zn-poor CZTS did not form in either of the synthesis conditions for C1 and C2. The fact that the CTS boundaries slightly cross the E- and D-type lines for Sn-rich compositions is unexpected. This crossing could be indicative of an overestimated Zn content in the samples when using EDS. The presence of the Cu3SnS4 phase is somewhat unusual. This phase, which can be written as Cu(I)2Cu(II)SnS4, requires Cu in its II oxidation state, and this is normally not seen under the conditions used here. We note that this same phase was also found in other combinatorial samples, which, as postulated by Lund et al.,18 might be as a result of different regions of the sample interacting via the gas phase. Cu3SnS4 does not naturally fit into the SPR model described above since there is no experimentally determined neutral defect that is compositionally equivalent to CZTS + Cu3SnS4. If this phase were to be incorporated in the model, we propose the hypothetical reaction below, in which under Zn-poor conditions, “dissolving” a small amount of Cu3SnS4 in CZTS results in formation of the neutral defect Cu(II)Zn in CZTS: Cu3SnS4 V 2Cu Cu + Cu(II)Zn + SnSn + 4SS i.e., Cu3SnS4 V Cu(II)Zn

(19)

The position of this equilibrium then defines the concentration of Cu(II) at the phase boundary toward Cu3SnS4. Since there is no gas phase component in this reaction, the position of the boundary should only depend on temperature. In agreement with this, the apparent phase boundaries in Figures 4a,b are essentially identical. For the Raman measurements at 532 nm, we initially detected only CZTS (and/or CTS), even for the Sn-rich regions where Sn-containing secondary phases are expected. Using further point measurements, it was determined that Sn−S phases were in fact present on the film surfaces, but as small and distributed flakes that were easily missed by the Raman spot in mapping. Moreover, in our process, Sn−S phases typically segregate to the back contact and thus could not be seen by surface Raman mapping.13,37 To analyze these phases, we scribed several additional lines across the samples, approximately perpendicular to the B−C lines, and scanned along these lines manually in the Raman system. Sn−S particles could be observed both at the CZTS surface and within the scribed lines (i.e., at the back contact). For C1, we could detect both SnS and SnS2, via the modes at 160−220 cm−1 and ca. 315 cm−1, respectively.38 As would be expected, these phases were only seen on the Sn-rich sides of the samples. Starting from the B−C (Cu/Sn = 2) line and moving into the Sn-rich side, SnS phases first became visible after some distance and grew continuously in size. Further from the B−C line, at around Cu/Sn = 1.75−1.80, intermixing of SnS2 was detected. Outside of the area mapped in Figures 3a and 4a, i.e., for Cu/Sn < 1.6, we could see 1−2 μm crystals of SnS2 on the sample’s top surface. In contrast to C1, sample C2 only showed SnS2 particles, at first only at the back contact, but for more Snrich compositions also on the surface. The observable SnS2 phases were far from the B−C line, at around Cu/Sn = 1.65− 1.70. Based on the observations, approximate phase boundaries for both SnS (C1 only) and SnS2 (both samples) could be determined.

Figure 4. Maps of the 785 nm Raman intensity in the range of the Cu3SnS4 main mode (316−327 cm−1) as a function of integral composition, for samples C1 (a) and C2 (b). The dotted line corresponds to the boundary between the presence and absence of observable Cu3SnS4 peaks in the spectra, from visual inspection. The intensity scale is in arbitrary units, and values cannot be compared from C1 to C2.

generated by integrating the spectra in the window 316−327 cm−1. For C1, the map shows a clear region of high intensity on the Zn-poor side. For C2, the same high-intensity region is present, but the contrast with the rest of the map is limited. This is because the CZTS modes in C2 are very intense, possibly due to higher ordering (see later). Therefore, for both samples we used visual inspection of the spectra to determine the presence/ 4631


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Chemistry of Materials Cu−S secondary phases were not seen by Raman spectroscopy at any of the chosen wavelengths, most likely due to an offstoichiometric Cu/S ratio, which can result in low Raman efficiency.39 Thus, identification of Cu−S phases was made by SEM imaging combined with EDS point measurements (at 10 keV accelerating voltage). From experience, we know that Cu−S phases segregate to the top surface of CZTS films during processing. Again, in line with expectations, Cu−S particles were only found on the Cu-rich sides of the samples (see Supporting Information). Analysis of the baseline sample C1 showed Cu−S particles with the approximate composition of Cu2S, with size variation from 0.25−1.5 μm in the Zn-rich, Cu-rich quadrant, and composition closer to CuS in the Zn-poor, Cu-rich quadrant. Although the S-content of the Cu−S phases appears to change, the apparent phase boundary is contiguous. In C2, a similar behavior was seen. In addition, these particles contained a slight contribution of Sn, and small ( 2, CZTS {i} complexes are present, which contain Cu or Zn interstitials. In both these situations, vacancy-assisted or interstitial-assisted exchange mechanisms will reduce the activation barrier for Cu−Zn exchange. This means that equilibrium ordering is more quickly reached during cooling, resulting in higher final QI value. Enhancement of atomic exchange (or more generally, migration) by the action of vacancies or interstitial defects is a well-understood process; see, for example, refs 8,43. The extent to which vacancies or interstitials can accelerate Cu−Zn exchange and the resulting QI value ought to depend upon their concentration in the lattice, i.e., on the CZTS phase composition.8 Due to this, we can use QI as an indicator of the relative Cu/Sn ratio in the CZTS phase, i.e., a higher value of QI for a given cooling rate indicates a higher concentration of vacancies (for Cu/Sn < 2) or interstitials (for Cu/Sn > 2). This is a useful tool since the EDS measurement gives the integral composition and cannot say anything about the partitioning of material between the CZTS and secondary phases. We note that in other reports, the additional intensity ratio Q′ has been proposed to distinguish certain CZTS types (A, B, and stoichiometric CZTS).29 However, our findings for the full compositional range show that Q and Q′ in fact convey the same information and can only be used to distinguish between the categories of defect type mentioned here, i.e., {s} and {v}/{i} types, on the basis of the enhanced ordering seen for the latter compared to the former.7

Figure 6. Maps of the secondary order parameter QI for the CZTS phase, determined from resonant Raman spectra for samples C1 (a) and C2 (b). The apparent secondary phase boundaries are included for orientation purposes, but their labels are omitted for clarity; see Figure 5 for details.

found in Part I (labels are omitted for clarity). In C1, there is an extremely uniform dependence of QI on the Cu/Sn ratio, with virtually no dependence on the Zn/(Cu + Sn) content. In the Cu-rich region, QI is large and constant. In the center of the map there is a band of low QI values spanning the range approximately 1.8 < Cu/Sn < 2.05. QI increases again at the very lowest Cu/Sn ratios, although not reaching the same level as on the Cu-rich side. In sample C2, the strongest trend is still in the Cu/Sn direction. While the Cu-rich side looks similar to C1, the degree of ordering on the Sn-rich side in C2 is dramatically increased. In some areas, QI reaches much higher values than observed anywhere on the baseline sample. The maximum QI values on the Sn-rich side reach ca. 4.6, in the region between the E- and Jtype lines, which is higher than for any CZTS sample we have produced using our baseline process. In C2, there is also some 4633


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Chemistry of Materials To more closely analyze the effect of Cu/Sn ratio on QI, “cross sections” of QI for a the range of Zn/(Cu + Sn) values greater than 0.34 are plotted for both samples in Figure 7. We analyze

behavior is exactly what is observed: In both C1 and C2, the QI parameters rise rapidly from Cu/Sn = 2 and saturate just above that value. The precise point of saturation cannot be accurately specified given the errors in the EDS measurement and in determination of QI, but it more or less coincides with the appearance of observable Cu−S phases at Cu/Sn ≈ 2.05−2.1. For the rest of the Cu-rich region, the QI values are constant; the CZTS phase is no longer changing. From this, we conclude that, on the side where Cu/Sn > 2, a small amount of Cu-excess can be included in the CZTS phase, which means that {i} complexes are present. Despite their low concentration, the interstitial defects in these complexes have a very strong effect on Cu−Zn ordering, resulting in QI values up to ca. 2.75, nearly three times higher than for CZTS with a stoichiometric Cu/Sn ratio. Since the limiting QI values on the Cu-rich sides were similar for both samples, we conclude that the partitioning of excess Cu between the CZTS phase and secondary phases was not affected by the anneal conditions. This is reasonable from the theory developed earlier: the solubility of {i} complexes should not depend on the concentrations of gas phase components since these complexes arise from equilibria involving only stable solid phases (Reactions 2, 6, 7, and 9). On the Sn-rich side of the samples, where Cu/Sn < 2, the behavior is more complicated. For C1, annealed with only sulfur vapor, QI is at first rather flat at its minimum value and only starts to rise below Cu/Sn ≈ 1.85. This rise coincides with the appearance of SnS secondary phases (Cu/Sn ≈ 1.87). Thus, although there was an “apparent” single phase region in the range of integral compositions from Cu/Sn = 1.87−2.0 (Figure 5a), the failure of QI to increase as integral Cu/Sn drops away from 2 suggests that the Cu/Sn ratio in the CZTS phase actually remained near-stoichiometric. Thus, the CZTS phase in this region must be either B- or C-type, or fully stoichiometric, depending on the Zn-content. The fact that the integral composition was nevertheless Sn-rich in this region can be due to several possibilities. One is that this region actually did contain Sn−S secondary phases, but that they were not observable by Raman due to being too small, or to being buried within the CZTS film. Another possibility is that the resonanceRaman measurement of QI is more surface-sensitive than the EDS measurement, and that depletion of Sn from the CZTS phase is strongest at the surface. This would result in a B/C-type surface, giving a low QI value, while the bulk of the film retained a Sn-rich composition (and may have been more ordered), explaining the Sn-rich integral composition. In either case, the conclusion is the same: under the synthesis conditions for C1, the CZTS phase on the Sn-rich side is unstable with respect to increased Sn-incorporation, and formation of {v} defect types is limited. For the baseline annealed sample, therefore, we conclude that the SPR is very small, or even nonexistent, toward the Sn rich side. The QI values do start to increase below Cu/Sn = 1.85, showing that, for sufficiently Sn-rich compositions, the Sn-content in the CZTS phase can be increased even in these annealing conditions (even at the surface). However, this is only occurring in the presence of detectable Sn−S phases, and even then, QI only reaches a maximum of around 2. The simultaneous increase in CZTS Sncontent and growth in the amount of Sn−S secondary phases is proof that these annealing conditions are far from equilibrium for the Sn-rich composition range. The behavior of C2, annealed in the presence of SnS2, is very different. This sample shows rapidly rising QI values as soon as Cu/Sn drops below 2. QI continues to rise at least until the

Figure 7. Profiles of secondary order parameter QI across the combinatorial samples C1 (a) and C2 (b), plotted against the integral Cu/Sn ratio. The shaded areas indicate where the labeled secondary phases were observable by Raman scattering, and the dashed lines indicates stoichiometry in the Cu/Sn ratio. The bold lines are guides to the eye.

the shape of these profiles on the basis of the above-described dependence of QI on the Cu/Sn content in the CZTS phase. One feature that is expected is that the smallest QI values should coincide with the composition ratio Cu/Sn = 2, where only {s} complexes are formed. This is in fact observed in both samples, within errors, which gives some confidence in the Cu/Sn ratios determined from EDS. Going away from Cu/Sn = 2 in either direction, assuming that the single phase region has some extension away from this point, the composition of the CZTS phase should change as the integral composition changes. The change should be continuous until the solubility limit of the relevant defect complexes is reached; {v} complexes for Cu/Sn < 2 and {i} complexes for Cu/Sn > 2. Beyond the solubility limits, the Cu/Sn ratio in the CZTS phase should be constant for a given Zn/(Cu + Sn) ratio, and any further increases in the integral Cu/Sn ratio will result in secondary phases appearing instead. Since the QI values are expected to reflect the stoichiometry of the CZTS phase only, they should vary only within a certain range of Cu/Sn around the stoichiometric value, i.e., only within the SPR, and they should be constant outside this range. On the Cu-rich side of the combinatorial samples, this 4634


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Chemistry of Materials observed SnS2 phase boundary is reached, at Cu/Sn ≈ 1.67. Typical values of QI at this boundary are in the range of 3, although as noted values up to 4.6 can be seen in the limited composition range near to Zn/(Cu + Sn) = 0.34. This is very close to the ideal behavior expected (and observed on the Curich sides in both samples), although it is unclear whether QI saturates or not due to the lack of data points for lower Cu/Sn ratios. It must be said that we still cannot rule out the presence of undetectable secondary phases in the region from Cu/Sn 1.67− 2.0. However, since the QI values are much higher in C2 compared to C1 for the same integral Cu/Sn ratio, the CZTS phase must be more Sn-enriched, and there must be correspondingly fewer (or no) secondary phases to be consistent with the integral composition. We conclude that the annealing conditions for C2 were capable of stabilizing CZTS with lower Cu/Sn ratios, in other words that the SPR toward Sn-rich compositions is larger. Again, the behavior on the Sn rich sides in C1 and C2 is wellexplained by the theory developed earlier. C2 was prepared using high partial pressures of S and SnS, due to the presence of excess solid SnS2 from the start. This is expected to stabilize {v} complexes, giving an extended SPR since the equilibrium positions in Reactions 12−15 will lie far to the right. However, C1 was prepared with only added S. Despite having Sn-rich integral compositions in the film, the equilibrium positions in the same Reactions 12−15 will now lie toward the left, resulting in the evaporation of S and the precipitation of SnS secondary phases, instead of the formation of {v} complexes in the CZTS phase. Thus, Sn-rich CZTS cannot be formed despite a Sn-rich integral composition, and the SPR should therefore have little or no extension toward Sn-rich compositions. So far, we have treated the two anneal experiments as if they provided steady-state conditions. This is actually not the case; in our baseline annealing process we have observed a strong effect of annealing time due to the loss of S partial pressure from the annealing zone as the anneal proceeds.44 This is also widely observed by other groups involved in thin-film CZTS(e) preparation, since it is inherently difficult to confine a gas with such a high boiling point as S (or Se) in a practical reactor with reasonable throughput. S or Se will tend to migrate to colder regions and condense, which acts as a “pump” to remove the gases from the reaction zone. In the study of Ren et al.,44 the stable Sn−S phase changed from SnS2 to SnS after about 3 min of annealing, in response to the gradual loss of S from the reactor. Thus, for an anneal of less than 3 min in the baseline conditions, we could expect a similar situation to sample C2: SnS2 stable in the solid phase, {v} complexes reaching the limiting concentrations dictated by Reactions 1, 5, 8, and 10, and the single phase region on the Sn-rich side having the same extent as in the equilibrium phase diagram. However, if the anneal continues for longer than 3 min, as for C1, SnS2 is no longer stable and starts to convert to SnS via 11. In this situation, the solubility limits of the {v} defect complexes are now governed by Reactions 12−15 and will drop as the S partial pressure goes down. If the CZTS phase was Sn-rich to start with, E-type complexes will now be lost via precipitation of SnS, Aand J-type complexes will be converted to B-type via the same process, and H-type complexes will be converted to C type. As this continues, the Cu/Sn ratio in the CZTS phase will get closer and closer to the stoichiometric value of 2, as found in B-type, Ctype, or fully stoichiometric CZTS. The film surface is most sensitive to the gas phase composition, so losses of SnS and S will occur there first, possibly resulting in composition gradients and

gradients in the dominant defect type throughout the depth of the film. Since the anneal atmosphere was continuously changing, sample C1 does not represent a given equilibrium situation, hence why the QI values kept changing outside the secondary phase boundaries, instead of saturating (i.e., the region with integral Cu/Sn ratio below ∼1.85 in Figure 7a). Figure 8a schematically illustrates the loss of E-type complexes

Figure 8. (a) Schematic of loss of an E-type defect complex from a Snrich CZTS surface, resulting in evaporation of SnS and S2 and a stoichiometric surface. (b) Limits of S2 partial pressure for stabilizing CZTS with various Cu/Sn ratios, for a saturated partial pressure of SnS.

from a CZTS surface, under conditions where a loss of SnS and S partial pressure reduces the solubility of the complex, with the result that the surface becomes stoichiometric. This corresponds to reaction 13 proceeding from right to left. Sample C2 was annealed with excess SnS2, and while most of this had been converted to greyish SnS by the end of the annealing period, the sample itself displayed only SnS2 secondary phases on its surface. Therefore, this sample was under quasi-steady-state conditions 4635


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as SnZn, present in parts of the Sn-rich composition range, may be detrimental.5,48 However, given the challenges involved in preparing genuinely Sn-rich CZTS (i.e., very high S2 pressure limits), in our view it is possible that we have not yet been able to fully explore this composition range of the CZTS SPR. Therefore, perhaps it is too early to conclude on the electronic properties of Sn-rich CZTS. It is to be hoped that, by extending the range of accessible composition of the CZTS phase, we can also extend the window of opportunity for defect engineering and tune the bulk properties of this material toward higher solar cell performance. Finally, while the experimental and discussion parts of this Article have only concerned sulfide CZTS, the principle of the model determining the SPR extent applies equally to the selenide and sulfo-selenide compounds. The important properties that could lead to differences in the observed SPR sizes in selenide vs sulfide CZTS are the formation energies of the defect complexes concerned and the vapor pressures of SnSe and Se.

and is at least close to being representative of the full, solid-state equilibrium situation that is captured in the equilibrium phase diagram.11 The model outlined at the start of this paper, Reactions 1−19, seems to correspond rather well to the experimental observations. We can therefore attempt to use this model going forward to help improve synthesis conditions for CZTS thin films for use in solar cells. First, it is worth pointing out the parallels between our defect equilibria and the decomposition reaction of the CZTS phase itself, which has also been shown to occur for too-low partial pressures of SnS and S2:19,45 Cu 2ZnSnS4 V Cu 2S + ZnS + SnS(s/g) + S2(g)

(20)

The minimum value of the S2 pressure needed for keeping this equilibrium toward the left side, given a saturated SnS partial pressure, is of the order of 10−4 mbar at 550 °C.19 This was determined from CZTS samples with an initially stoichiometric integral Cu/Sn ratio, by finding the steady-state S2/SnS partial pressures at which there was no net change in Cu/Sn ratio upon extended (5 h) annealing. In the light of our defect equilibria model, we learn that this pressure limit is not universal, but is only applicable to the case of CZTS with Cu/Sn ≥ 2. If we wish to make CZTS with a lower Cu/Sn ratio, i.e., with a higher concentration of {v} complexes (A-, E-, H-, or J-type), the S2 pressure limits will be higher. In principle, one could determine analogous S2 pressure limits for stabilizing CZTS with different concentrations of {v} complexes. While we do not do this here, we can suggest that the S2 pressure limit for the maximal concentration of {v} complexes, i.e., the maximal extension of the SPR toward Sn-rich compositions, ought to be the same as the limit for stabilization of SnS2. If that is so, and given a saturated SnS partial pressure (2 × 10−3 mbar at 550 °C), the S2 pressure needed to obtain a CZTS phase with the highest possible concentration of {v} complexes should be around 1 mbar at 550 °C,19 i.e., 4 orders of magnitude higher than the S2 pressure needed to stabilize CZTS with a stoichiometric Cu/Sn ratio. Figure 8b shows the pressure limits for stability based on this argument, for a range of temperatures. Thus, if vacancy-rich A-, E-, H-, or J-type CZTS is of interest for solar cells, it is highly recommended that synthesis processes and reactors should be re-evaluated to see what S2 pressures they can actually provide. A simple way to do this, as we showed recently,44 is to attempt to deposit and/or anneal SnS2 (powder or thin film). If this can be done successfully without decomposition to SnS, then the full range of the CZTS SPR can also be probed in the given experimental approach. As an alternative to increasing the S2 pressure, adding a Sn-bearing component to the thermal step in synthesis, such as SnS, SnS2, or SnCl2 could help. Alternatively, very short reaction times could be used in scenarios where there is loss of partial pressure from the reactor during processing. One reason that the vacancy-rich/Sn-rich region of the CZTS SPR could be interesting is that it allows for considerably lower levels of Cu−Zn disorder (as seen here and in ref 7). If it is indeed the case that disorder contributes to the voltage deficit in CZTS solar cells, Sn-rich compositions could be one means of reducing the problem. Whether or not it is related to ordering, there are indications of a beneficial effect of high Sn-content in the CZTS phase,44,46,47 and it is known that the highest device efficiencies for CZTS have resulted from very short anneals (see ref 3 and references therein), which might indicate the need to preserve very high S2/SnS partial pressures. However, some authors have suggested that certain Sn-related point defects such

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CONCLUSIONS Cu2ZnSnS4 (CZTS) is hoped to be a future, earth-abundant absorber material for thin film solar cells, but recently, progress in device performance has saturated below the level that is needed for commercialization. In this work, the question of the size of the single phase region of (CZTS) obtained from thin film synthesis methods was addressed, with a view to determining the potential for defect engineering and thereby future improvements in solar cell performance. A chemical model for the single phase region was presented, based on equilibria between the defect complexes responsible for offstoichiometry in the CZTS phase and the other solid- and gasphase components present during film synthesis. The model predicts a variable single phase region size, depending on the partial pressures of SnS and S2. In particular, the possibility to form Sn-rich CZTS, containing so-called A-, E-, H-, or J-type defect complexes (more generally, containing Cu or Zn vacancies) is predicted to be reduced for too-low partial pressures of these gas phase species. The model was evaluated by analysis of combinatorial thin-film CZTS samples prepared with different conditions. Raman and SEM-EDS characterization of secondary phases allowed the outer limits of the single phase regions to be mapped. The Cu/Sn ratio in the CZTS phase could be qualitatively determined by analysis of the secondary order parameter from resonant Raman, thanks to the influence of vacancies and interstitial defects, present in different composition ranges, on the ordering kinetics. The predicted behavior was observed, with a substantially higher ordering, linked to a higher Sn-content, being obtained in the CZTS phase for a film prepared with higher partial pressure of SnS. CZTS annealed according to “standard” conditions, however, had an extremely narrow or possibly nonexistent single phase region toward Sn-rich compositions. The single phase region toward Cu-rich compositions was not affected, again in line with predictions from our model. The important implication from our findings is that typical synthesis strategies are not capable of accessing the full range of the CZTS single phase region, with the result that our understanding of CZTS defect chemistry from experimental studies is still incomplete, especially for Snrich compositions. There is evidence that improved synthesis approaches, based on these results, could lead to better control of the CZTS stoichiometry, allowing more effective defect engineering and, possibly, more favorable optoelectronic properties. Synthesis conditions to access the full range of the 4636


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single phase region were mapped out, and suggestions were made about how to achieve these in practice.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.8b01213. Example Raman spectra for the phases analyzed in this study and further details on Sn−S and Cu−S phase identification, as well as comparative X-ray diffraction measurements on sample C2 (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Jonathan J. S. Scragg: 0000-0001-8686-8721 Author Contributions

All authors have given approval to the final version of the manuscript. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors gratefully acknowledge funding from the Swedish Research Council, The Swedish Foundation for Strategic Research, and STandUP for Energy.

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REFERENCES

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The single phase region in Cu2ZnSnS4 thin films from theory and

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