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# Vapor Deposition of Hybrid Halide Perovskites in Ultra-High Vacuum: Challenges Revealed by Probing the Organization Dynamics in Real Time

Abdel R Allouche, Christian P Sonny Tsotezem, Arindam Mukherjee, Elena M Staicu Casagrande, Anouchah Momeni, Aimeric Ouvrard, Yimin Guan, Laurent Guillemot, Daniela Torres-Díaz, Lionel Amiaud, et al.

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## ► To cite this version:

Abdel R Allouche, Christian P Sonny Tsotezem, Arindam Mukherjee, Elena M Staicu Casagrande, Anouchah Momeni, et al. Vapor Deposition of Hybrid Halide Perovskites in Ultra-High Vacuum: Challenges Revealed by Probing the Organization Dynamics in Real Time. Journal of Physical Chemistry C, 2025, 129 (41), pp.18737-18748. 10.1021/acs.jpcc.5c04748. [hal-05379707](https://hal.science/hal-05379707v1)

HAL Id: hal-05379707

https://hal.science/hal-05379707v1

Submitted on 9 Dec 2025

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Vapor deposition of Hybrid Halide Perovskites in

Ultra-High Vacuum: challenges revealed by probing

the organization dynamics in real-time.

Abdel R. Allouche<sup>1,2,*</sup>, Christian P. Sonny Tsotezem<sup>1</sup>, Arindam Mukherjee<sup>1</sup>, Elena M.

Staicu Casagrande<sup>1</sup>, Anouchah Momeni<sup>1,3</sup>, Aimeric Ouvrard<sup>1</sup>, Yimin Guan<sup>1</sup>, Laurent

Guillemot<sup>1</sup>, Daniela Torres-Díaz<sup>1</sup>, Lionel Amiaud<sup>1</sup>, Anne Lafosse<sup>1</sup>, Adnan Naja<sup>2</sup>, Gaëlle

Trippé-Allard<sup>4</sup>, Emmanuelle Deleporte<sup>4</sup>, Hocine Khemliche<sup>1,*</sup>

<sup>1</sup> Université Paris Saclay, CNRS, Institut des Sciences Moléculaires d'Orsay, 91405 Orsay, France.

<sup>2</sup> Laboratoire de Physique et Modélisation (LPM), EDST, Lebanese University, 1300 Tripoli, Lebanon

<sup>3</sup> CY Cergy Paris Université, 95000 Cergy, France

<sup>4</sup> Laboratoire Lumière, Matière et Interfaces (LuMIn), Université Paris-Saclay, ENS Paris-Saclay,
CentraleSupélec, CNRS, 91190 Gif-sur-Yvette, France

Abstract

With record efficiencies and reduced costs, hybrid halide perovskite solar cells could revolutionize photovoltaics, provided that the crippling instability problem is resolved. To reveal the relationship between structural defects and instability, vacuum deposition emerges as a promising alternative to the prevalent solution method. In addition to responding to the major upscaling issue, the vacuum approach offers a greater control over the film growth. Vacuum evaporation could also help elucidating the puzzle regarding the link between morphology and optoelectronic properties. Paradoxically, layers prepared in solution, polycrystalline by nature and containing numerous defects, have so far achieved better performances than layers prepared in vacuum, although the latter show better crystalline
quality. Here, we describe an original approach, setting the conditions for the best possible control of
the layer growth. Deposition is performed in ultra-high vacuum (UHV) and the growth dynamics
characterized in real-time by Grazing Incidence Fast Atom Diffraction (GIFAD). Applied to the
prototypical system CH3NH3Pbl3, we demonstrate that chemical reaction between the co-evaporated
precursors, CH3NH3I and Pbl2, cannot take place at room temperature. In the sequential mode, pure
Pbl2 layers of high crystalline quality cannot be converted to CH3NH3Pbl3 by subsequent evaporation
of CH3NH3I. These findings raise fundamental questions regarding the role of contaminants from the
residual gas when CH3NH3Pbl3 layers are deposited in high vacuum conditions with unbaked vessels.
These results therefore also question recent molecular dynamics models describing the reaction steps
between Pbl2 layers and CH3NH3I molecules that intercalates in the van der Waals space of the
inorganic network.

Keywords

Hybrid Perovskite – Thin Films - Growth Dynamics – Fast Atom Diffraction – Real-time characterization
– Molecular Beam Epitaxy

TOC Graphic

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Introduction

Hybrid halide perovskites have shown great promise for highly efficient and cost-effective optoelectronics devices. Used as active layers in solar cells, these materials have now reached a record 27% Power Conversion Efficiency (PCE)¹, similar to that of crystalline silicon. Despite significant research efforts, many challenges remain that include solving the stability issue² and its relation to defects, identifying the nature and influence of the interfaces³ and achieving a better understanding of their crystallization dynamics. Solution-processed layers are polycrystalline; the grain boundaries are responsible for the rapid structural degradation by reaction with water⁴; ion migration through vacancies also contributes to degradation⁵. It is therefore expected that higher crystallinity could partially resolve the stability issue.

Vacuum deposition could be a good alternative to the widely used solution method. Indeed, since the initial work by the H. Snaith's group⁶, many studies reported a higher structural quality, a smoother surface, and a more homogeneous thickness for vacuum processed layers⁷⁻⁹. However, very early, drawbacks of the vacuum deposition method were identified. The high volatility of the organic compound makes a well-balanced stoichiometry difficult to achieve¹⁰, while the interface with the substrate is almost inevitably composed of PbI₂, which has drastic consequences on transport properties¹¹.

To extract the best from vacuum evaporation, we describe a novel approach in which the growth is performed in ultra-high vacuum (UHV) and its dynamics characterized in real-time by Grazing Incidence Fast Atom diffraction (GIFAD). GIFAD is based on the coherent scattering of keV energy He atoms at angles typically below 1° with respect to the surface plane¹²,¹³. In this geometry, the He atoms are reflected by the electron density above the last atomic plane, resulting in a very smooth interaction and extreme surface sensitivity¹⁴. The advantages of GIFAD lie in its compatibility with insulators¹⁵ and fragile materials such as organics¹⁶, in addition to the quantitative determination, from the diffraction pattern, of the equipotential energy surface¹⁷. To a first approximation, the latter can be viewed as the
electron density profile measured at typical distances of 2-4 Å above the last atomic plane, i.e., for distances that are the most meaningful of gas-surface interactions.

This work focusses on the early stage of the growth, during the simultaneous or sequential evaporation of CH₃NH₃I (MAI) and Pbl₂ on Ag(001). The objective is to resolve the organization dynamics leading to crystallization on the atomic scale and from the very first adsorbed molecules. For this reason, a single crystal substrate exhibiting a GIFAD pattern is a prerequisite, thus allowing a clear insight onto the growing layer properties such as orientation and commensurability. To complement the information provided by operating GIFAD in real-time, spontaneously desorbing species are collected by a Quadrupole Mass Spectrometer (QMS). The latter is also used for chemical analysis during thermal desorption of the deposited layer. Although GIFAD provides all the crystallographic information relative to the top layer, as well as direct signatures of coexisting domains and mosaicity, Low Energy Electron Diffraction (LEED) can be performed *in situ* to confirm the GIFAD data. *Ex situ* characterization includes High Resolution Electron Energy Loss Spectroscopy (HREELS) for accessing vibrational properties in the range 0–4000 cm⁻¹ ¹⁸ by scattering low energy electrons, X-Ray Photoelectron spectroscopy (XPS) to quantify the chemical composition of the grown layer, and Scanning Tunneling Microscopy (STM) for probing local morphology. The results demonstrate that, on top of an inevitable and self-terminated MAI buffer layer, co-evaporation of Pbl₂ and MAI inevitably produces a Pbl₂ film with high crystalline quality with no organic content. Nonetheless, we do observe a remarkable dynamics at the very first stage of Pbl₂ adsorption, pointing to a chemical reaction with the buffer layer.
This translates into a displacive phase transition in the surface plane, from the 2D square phase of the MAI buffer layer to a hexagonal Pbl₂-rich termination layer.
# 1. Experimental Section

The experimental approach combines several surface sensitive techniques that operate either in real-time during layer growth (GIFAD), in situ for post-growth characterization (LEED) or ex-situ following sample transfer (HREELS, STM, XPS).

**2.1. GIFAD.** The probe beam is prepared as follows: He⁺ ions are extracted from an ion source at energies typically in the range of 500–1000 eV. Following neutralization in a He gas cell; a nearly parallel neutral beam is produced by a set of collimating apertures with a size lower than 0.2 mm and directed to the sample surface at grazing angles that can vary between 0.25 and 1.0°. Reflected atoms are collected on a position sensitive detector made of a set of two Microchannel plates and a phosphor screen, and images captured by a CCD camera.

**2.2. Substrate preparation.** The Ag(001) substrate was prepared very carefully by cycles of ion sputtering at 1 keV (initial cycles) and 500 eV (final cycles) and annealing at 500 °C until a well-contrasted diffraction pattern is observed along the [110] direction. Along the [100] direction, only the specular peak (0th order) could be observed due to the very small projected corrugation in this direction.

**2.3. Evaporation.** MAI and Pbl₂ are introduced in glass crucibles that are radiative-heated inside a three-cell UHV evaporator. The crucibles are degassed for hours at temperatures (80 °C and 270 °C for MAI and Pbl₂ respectively) close to the evaporation values; a cooling system prevents temperature crosstalk between the two crucibles.

Pbl₂ (99.999% - ultradry) was purchased from Thermo Fisher and used as received. MAI was synthesized by adding methylamine (33% in ethanol) (25 mL, 200 mmol) dropwise to a cold solution of iodohydric acid (57% in water) (28 mL, 1.05 eq., 210 mmol) in ethanol (500 mL). The mixture was stirred for 2 hours at 2 °C. The solvent was removed until the first crystals appeared. After cooling to 4 °C, the crystals were filtered and recrystallized in ethanol. The white MAI crystals were then washed with diethyl ether and dried overnight at 60 °C.
**2.4. HREELS.** The HREEL spectrometer is housed in a separate UHV chamber equipped with a load-lock system to facilitate sample transfer. The spectrometer consists of a double monochromator and a single analyzer (model IB 500 by Omicron). The spectrum was recorded in the specular geometry ($\theta_i = \theta_f = 55^\circ$ with respect to the surface), at an incident energy of 6 eV. Data acquisition is performed at 23 K in order to improve the resolution, the latter shows a Full Width at Half Maximum (FWHM) of 3.5 meV measured on the elastic peak.

**2.5. STM.** The STM, operating in UHV at room temperature (ScientaOmicron) uses tungsten tips prepared by electrochemical etching in KOH. STM images were processed using WSxM software¹⁹.

**2.6. XPS.** The XPS measurements are performed in a separate UHV chamber with a load-lock allowing sample transfers with a UHV suitcase. In our non-monochromatised, dual-anode Xray source from PSP Vacuum, we use the Mg anode, providing a main line $K\alpha = 1253.6$ eV and a hemispherical electrostatic analyzer RESOLVE 120 from PSP Vacuum. Wide energy scan spectra are acquired with the CRR10 mode (Constant Retard Ratio = 10), while the narrow energy spectra are recorded using the CPE20 mode (Constant Path Energy = 20 V) with a 0.8 eV energy resolution.

## 2. Results

Difficulties to control the deposition rate of MAI, which derive from its high volatility, are usually resolved by controlling either its partial pressure or the total pressure in the vacuum chamber during the deposition. This is however not sufficient since the sticking coefficient may vary during the different stages of the deposition. In UHV, we have noticed that MAI spontaneously dissociates in the crucible even at room temperature into gaseous sub-products, $CH_3NH_3$ and HI, whose partial pressures remain in the $10^{-11}-10^{-12}$ mbar range depending on process history. Despite these very small concentrations, these sub-products eventually produce a complete buffer layer on the Ag(001) substrate over 12 to 24 hours. Several days after the last deposition process, MAI dissociation products are still present in the
residual gas (see the mass spectrum in Figure S1 on the Supporting Information). The major fragments
are CH<sub>3</sub>NH, CH<sub>3</sub>NH<sub>2</sub><sup>+</sup>, CH<sub>3</sub>NH<sub>3</sub><sup>+</sup>, I<sup>2+</sup>, I<sup>+</sup>, HI<sup>+</sup>, and CH<sub>3</sub>NH<sub>3</sub>I, corresponding to m/z values of 30, 31, 32, 63.5,
127, 128, and 254, respectively. These peaks were similarly observed in the work of Juarez-Perez and
his team<sup>20</sup>.

Real-time monitoring by GIFAD shows that this self-terminated buffer layer, one monolayer (ML) thick, inevitably forms as soon as the substrate temperature is lower than 420°C, and it is strictly identical whether it forms by the slow deposition or directly by thermal evaporation from the crucible. This observation confirms the study by Nakamura et al.<sup>21</sup> who demonstrated that two independent gas sources (CH<sub>3</sub>NH<sub>3</sub> and HI) produce the same MAI layer as with a single solid CH<sub>3</sub>NH<sub>3</sub>I source.

Our growth procedure starts with the controlled deposition of the buffer layer, the latter exhibits a c(2x2) structure similar to that obtained by depositing iodine on silver<sup>22</sup>. GIFAD and LEED data are presented in Figure S2 on the Supporting Information. A thorough analysis of the buffer layer properties will be reported in a future publication. Here, we are primarily concerned by its chemical composition since one would expect the organic species to desorb spontaneously, due the known affinity between iodine and silver. Chemical species can be identified by their vibrational properties; for this purpose, the freshly prepared sample is rapidly transferred in air to the HREELS setup for further analysis. In this UHV setup, the sample is first cooled down to 23 K, then rapidly heated up to 220 K at a rate of 40 K.min<sup>-1</sup> to remove the residual water that accumulates while cooling down the cryostat, and finally cooled down again to 23 K before data acquisition. The energy loss spectrum of Figure 1 shows well-resolved structures listed in Table S1 of the Supporting Information, together with their attribution according to published data from different vibrational spectroscopic techniques<sup>23,24</sup>. The general shape of the structured band extending from 90 up to 180 meV, which results from the contributions of many overlapping modes, matches well with the pattern obtained on a pristine MAPbI<sub>3</sub> layer by electron energy loss spectroscopy in a transmission electron microscope (TEM-EELS)<sup>25</sup>.
In Figure 1, the methyl group (-CH₃) gives rise in particular to the broad loss band centered around 361 meV (stretching modes) and to the losses resolved at 181 and 170 meV (bending modes). The -NH₃⁺ group contributes as shoulder features at ~386 meV (stretching modes) and ~196 meV (bending modes). Note that NHₓ stretching modes are expected to have low intensities in HREELS²⁶. The C-N bond contributes through its resolved torsional (twist along CN axis) mode at 30 meV and its non-resolved stretching mode v(C-N) around 118 meV. We conclude that the buffer layer obtained by evaporating MAI on top of Ag(001), effectively retains the MA species.

<figure><img src="image_4.png" /><figcaption>**Figure 1.** HREEL spectrum of the MAI buffer layer on Ag(100) measured at 23 K in the specular geometry after annealing at 220 K to remove residual water adsorption (Incident energy: E<sub>HREELS</sub> = 6 eV, ΔE<sub>FWHM</sub> = 3.6 meV, Elastic peak intensity: I<sub>0</sub> = 7845 cts/s, Incidence and exist angles: θ<sub>i</sub> = θ<sub>f</sub> = 55°).</figcaption></figure>

Noticeably, the peak at 30 meV or 242 cm⁻¹ (M-A torsion mode) has been observed by Raman spectroscopy at 275 cm⁻¹ on MAPbI₃ but not on pure MAI²⁷. The study by Niemann and co-workers²⁸ demonstrates that the M-A torsion frequency strongly depends on the halogen species in CH₃NH₃PbX₃ (with X = Cl, Br, I); the heavier the halogen, the stronger the shift to lower frequencies. The authors interpret these observations as due to the strong MA interaction with the surrounding halogens,
resulting in the MA deformation. The very low frequency measured here for the MA torsion could
suggest that it is strongly interacting with the I-Ag interface. The achieved resolution and the sharpness
of some of the observed losses (e.g. at 30, 161, 169 and 181 meV) indicate that the high crystalline
quality of the buffer layer, as characterized by GIFAD following the deposition, survives the transfer in
air. For the sake of simplicity, we will refer to this layer as the MAI buffer layer.

Growth mode and structural properties are simultaneously accessible by GIFAD. Figure 2 displays the
time evolution of the scattered intensity, along the diffraction direction, during a growth sequence
where we first evaporate MAI alone to produce a high-quality MAI buffer layer onto a previously
prepared Ag(001) substrate. The probe beam is aligned along the substrate [110] direction; the origin
of time is set at the start of the co-evaporation. For this sequence, we selected slow deposition rates
in order to resolve the properties of the interface as well as the various phase transitions. Upon
completion of the MAI buffer layer, we close the MAI shutter and let it stabilize for 700 s. Then, we
open both MAI and Pbl₂ shutters, at crucible temperatures of 80 °C and 250 °C, respectively. As can
be seen in Figure 2, the intensity distribution within the diffraction peaks evolves nearly
instantaneously. It follows a complex transition, whose details will be reviewed in the Discussion
section, that gradually transforms the square phase into new phase that only shows the zeroth order.
At t = 4300 s, the GIFAD reflectivity curve (see Figure S3) reaches a maximum, indicating completion of

<figure><img src="image_5.png" /><figcaption>**Figure 2.** Time evolution of the diffraction pattern along the Ag [110] direction, acquired with 800 eV He at 90 meV normal energy. The initial deposition of MAI is followed by co-evaporation of MAI and Pbl₂. A 15° azimuthal rotation is performed at t = 4300s, showing a new diffraction pattern. The left-hand side displays the Ag substrate.</figcaption></figure>
the new layer; a sample rotation of 15° around the azimuth leads to appearance of a new and well resolved diffraction pattern. Overall, the sequence produces exactly one and a half monolayer, as explained in section 4 of the Supporting Information.

An azimuthal scan, obtained by rotating the sample around its normal, performed after the growth sequence, shows that the final phase has a 12-fold symmetry, with nearly identical diffraction patterns repeating every 30°, as shown in Figure 3. Keeping in mind the 4-fold symmetry of the MAI buffer layer, the observed 12-fold symmetry is consistent with the coexistence of two hexagonal mirror domains

<figure><img src="image_6.png" /><figcaption>**Figure 3.** Diffraction patterns observed from the hexagonal phase with GIFAD parameters comparable to those of figure 2: (a) along the Ag [100] direction, labeled 0°. (b) Along the 30° direction. In both cases, the pattern shows an abnormally intense zeroth order contribution.</figcaption></figure>

rotated by 30° (or equivalently by 90°). This hypothesis is further supported by the abnormally high relative intensity of the zeroth order (Figures 3). An additional clue is provided by combing images acquired at different incidence angles (polar scan). The result is shown in Figure 4a, the distortion appearing at large angle (black rectangle) reveals that the diffraction patterns result from the superposition of two distinct patterns corresponding to the two main crystallographic directions of a hexagonal lattice, as observed by GIFAD on graphene²⁹. As shown in Figure 4b, we succeeded in fitting the whole diffraction pattern with two contributions, each described by a series of equally spaced pseudo-Voight peaks. One corresponds to a highly corrugated lattice direction (diffraction pattern with a large number of Bragg peaks) with a reciprocal vector of 1.49 Å⁻¹ and another with only orders up to ±2 (small corrugation) and a reciprocal vector of 2.58 Å⁻¹. The ratio between these two vectors is √3.
From considerations provided in the Supporting Information (Figure S4), we conclude unambiguously
in favor of a hexagonal lattice. We therefore identify the two structures in Figure 4b as corresponding
respectively to the armchair and zigzag directions. Note that the intensity distribution within the two

<figure><img src="image_7.png" /><figcaption>**Figure 4.** a) Diffraction chart built form a polar scan at 500 eV beam energy along the Ag [100] direction of the co-evaporated layer. The anomaly enclosed in the black rectangle reveals a second set of diffraction peaks. (b) A fit of the experimental diffraction spectrum from Figure 3a reveals two distinct contributions with drastically different corrugations (see Figure S4).</figcaption></figure>

zeroth orders is random since the relative weight between the mirror domains explored by the incident
beam is unknown, although the two types of domains are expected to have equal share on the whole
sample.

The coexistence of two hexagonal domains rotated by 30° with respect to each other, as represented
in Figure 5a, is confirmed by the *in-situ* LEED diffraction pattern shown in Figure 5b. The raw LEED
image has not been corrected for the distortions due to the LEED being equipped with a microchannel
plate (MCP-LEED). It is noteworthy that the contribution from the Ag substrate is clearly visible on the
LEED image (green distorted square), while that from the c(2x2) is hardly visible. This point will be
further considered in the Discussion section. To complement these macroscopic-scale structural
properties with a nanoscale perspective, the sample is then transferred to the STM apparatus within
a UHV suitcase; the images (Figures 5c and 5d), acquired in UHV and room temperature, also confirm
the existence of the two mirror domains. We observe that the MAI buffer layer is fully covered by the
co-evaporated layer. Hexagonal islands, 3 Å high, are found on top of this layer (inset of Figure 5d).
This information agrees with the coverage of exactly one and a half monolayer determined by the
GIFAD reflectivity curve (Figure S3) associated with the deposition sequence of Figure 2.

The crystallographic properties of the hexagonal lattice obtained at the end of the growth sequence
depicted in Figure 2 can now be quantified rigorously. Along both the 0° and 90° directions, the highly
structured diffraction pattern (blue spectrum in Figure 4b) yields a period of 4.20 Å, for a value of 4.04
Å for the 30° and 60° directions; these values need to be compared to the value of 4.09 Å for the lattice
parameter of silver. In agreement with GIFAD, STM shows that one of the two unit vectors ($\|\vec{a}_1\|=$4.6
Å) is 5% larger than the other one ($\|\vec{a}_2\|=$4.4 Å). These larger values derived from the STM images, in
comparison to those provided by GIFAD, are refined by recalibrating the STM images on the bare

<figure><img src="image_8.png" /><figcaption>**Figure 5.** a) Coexistence of mirror domains with hexagonal lattice. b) LEED diffraction pattern showing the two hexagonal lattices in addition to the square lattice of the substrate, the left inset shows the bare image; see text for the c(2x2) contribution. In a) and b), the Ag(001) lattice is displayed in green, the c(2x2) structure in white and the hexagonal domains in pink and yellow. c-d) STM images (I=15 pA, U= -0.7 V) showing the two perpendicular domains. The inset in d) shows the topography of a hexagonal island on which image d) was obtained. The region outside this hexagonal island provided the image in e). f) Topography line profiles from c, d and e).</figcaption></figure>
Ag(001) substrate. This procedure yields values, $||\vec{a}_1||=4.28$ Å and $||\vec{a}_2||=4.02$ Å, very close to those obtained by GIFAD.

We conclude that an anisotropic epitaxial constraint produces a strained irregular hexagonal lattice. The epitaxial character of the growth is clearly visible in Figure 6, when the hexagonal layer is analyzed by GIFAD along the Ag [100] direction. In this case, the diffraction peaks of the hexagonal lattice are well-aligned with those of the MAI buffer layer, the latter itself growing epitaxially on the silver substrate. Note that, due to the irregular shape of the hexagonal lattice, the actual angles between diffracting directions, generically labelled 0°, 30°, 60°, etc. should slightly deviate from these values.

<figure><img src="image_9.png" /><figcaption>**Figure 6.** Diffraction profiles from the MAI buffer layer and the co-evaporated layer vs transfered momentum (Å⁻¹) with the probe beam parallel to the Ag[100] direction.</figcaption></figure>

Our accuracy on the determination of the azimuthal angle, mainly due to mechanical play, has not been challenged for determining more precisely these angular directions.

The crystallographic properties being determined, we now concentrate on the layer content. The measured periods of this hexagonal layer (a = 4.20 Å and b = 4.04 Å) are very close to those of bulk Pbl₂, for which we expect a period of 3.96 Å in the [100] direction, for a lattice parameter of 4.56 Å (JCPDS N° 07-0235). This is significantly different from the parameters of the hexagonal MAPbI₃ predicted by E. Mosconi et al.³⁰. In the latter study, the calculated MAPbI₃ hexagonal phase, with a =
8.49 Å and b = 8.50 Å, exhibits a much greater stability, in comparison to the three well-known phases of MAPbI$_3$ (cubic, tetragonal, orthorhombic). We conclude in favor of a PbI$_2$ rich layer, suggesting that MAI does not co-adsorb with PbI$_2$.

To confirm this point, we repeated the growth sequence depicted in Figure 2 by evaporating PbI$_2$ alone on the MAI buffer layer; the resulting layer properties are remarkably similar to those described above. To eliminate any possible substrate effect at these short distances, we produced thicker layers in both modes, co-evaporation and evaporation of PbI$_2$ alone, in otherwise identical conditions. The respective data are compared in Figure 7. The reflectivity curves (Figure 7a), although comparable, show a weaker loss of intensity relative to the first layer for the co-evaporation, meaning that the final surface is smoother in this case. We also observe a slightly faster growth rate for the co-evaporation. This is quantified in Figure 7b; co-evaporation results in a growth rate that is 6.8% faster than for deposition of PbI$_2$ alone, though the properties of the final layers are identical in both crystallographic properties and intensity distribution within the diffraction peaks. This data suggests that during co-evaporation the impinging MAI molecules do have a noticeable influence on the growth mode of PbI$_2$, even though they eventually mostly desorb from the surface. A substance that improves the growth of a material and improves its surface quality is usually referred to as a surfactant. However, the terminology has somehow evolved to characterize a floating monolayer, of a material different from

<figure><img src="image_10.png" /><figcaption>**Figure 7.** Comparison between co-evaporation and deposition of PbI$_2$ alone; (a) GIFAD Reflectivity and (b) Deposition rate measured with a 800 eV He beam.</figcaption></figure>
both the substrate and the growing layer, that promotes the layer-by-layer growth of semiconductors³¹.

With a thickness of 11 ML, the influence of the substrate, which translated into an epitaxial constrain described earlier, is expected to diminish. The layer number dependence of the period seen by GIFAD is displayed in Figure 8, for the co-evaporation and the deposition of Pbl₂ alone. In both cases, we notice a full relaxation of the lattice parameter to a mean value of 4.05 Å at a thickness between 8 and 10 ML. This period corresponds to a lattice parameter of 4.67 c, very close to the reported bulk value (4.56 Å) of pure Pbl₂. Moreover, the same period is now measured along all directions, meaning that the lattice has turned into a regular hexagon. Nevertheless, a substrate effect may still be present and affect the chemical reactivity between MAI and Pbl₂ for the layer thickness investigated here. Such a substrate effect acting at larges distances has been observed in several studies on the vacuum growth of MAPbI₃.

<figure><img src="image_11.png" /><figcaption>**Figure 8.** Period seen by GIFAD along the substrate [110] direction the deposition of Pbl₂ alone and the co-evaporation mode, same GIFAD data as those used for Figure 7.</figcaption></figure>

In the work of Olthof and co-workers³², depending on the substrate (PEIE, ITO, PEDOT:PSS, and MoO₃), the stoichiometric phase only appears above a thickness of 3 to 30 nm, while the interface is rich in organic dissociation products. Conversely, Xu et al.¹¹ observed a Pbl₂ rich interface on various substrates (ITO, pedot/ITO, Si and glass). On ITO in particular, the I/Pb ratio increases monotonously
from ~2 at 5 nm thickness and reaches the value of 3, expected for stoichiometric MAPbI$_{3}$, only at 50 nm.

In the light of these studies, we deposited a much thicker layer in the co-evaporation mode, directly on the 11 ML thick film described in Figure 7. For this thicker film, the evaporation temperatures were increased from 100 °C to 128 °C and from 270 °C to 307 °C for MAI and PbI$_{2}$ respectively. The resulting deposition rate reaches 0.56 ML/min for PbI$_{2}$. Figure 9 displays the raw GIFAD reflectivity curve (top) and the result after a smoothing process (bottom); the curve shows 39 oscillations, which corresponds to an overall film thickness of 50 ML. The final layer is characterized by a hexagonal structure, with a unique diffraction pattern appearing at 30° intervals. As shown in Figure S5 on the Supporting information, the diffraction patterns are similar to those observed for the 11 ML thick layer (Figure 3), suggesting again a pure PbI$_{2}$ phase.

<figure><img src="image_12.png" /><figcaption>**Figure 9.** Deposition of 39 MLs by co-evaporation on top of the 11 MLs deposited previously. (a) raw GIFAD reflectivity, (b) smoothed GIFAD reflectivity that enhances the oscillatory behavior.</figcaption></figure>
To confirm this information, the chemical composition was verified by ex situ XPS measurements. The sample was transferred using a UHV suitcase, ensuring a clean vacuum environment. Figure 10 shows a wide energy scale spectrum of the emitted photoelectrons as a function of their binding energy. Several intense peaks can be identified as related to lead and iodine. The main contributions are the I(3d<sub>5/2</sub>) and I(3d<sub>3/2</sub>) at 619 eV and 631 eV, I(4d<sub>3/2</sub>) and I(4d<sub>5/2</sub>) close to 50 eV, as well as Pb(4f<sub>7/2</sub>) and Pb(4f<sub>5/2</sub>) at 137 eV and 142 eV; Pb(4d<sub>5/2</sub>) and Pb(4d<sub>3/2</sub>) are visible at 412 eV and 43415 eV. Meanwhile, it is worth noting that we do not observe any silver-related peaks; this is due to the film thickness being much larger than the electron elastic mean free path. Neither do we see any obvious signatures of

<figure><img src="image_13.png" /><figcaption>**Figure 10.** Full XPS spectrum displaying main lines at binding energies related to Pb and I. Positions of missing Ag, C, and N lines are shown.</figcaption></figure>

nitrogen or carbon atoms, pointing to the absence of organic species in the layer. Figure 11 shows a set of narrow scale high-resolution spectra. Figure 13a and 13b show the main Iodine and Lead contributions. The spectra were deconvoluted and background-subtracted. The I(3d<sub>5/2</sub>) peak is centered at the energy (619.5 eV) typically measured on MAPbI<sub>3</sub>,<sup>31</sup> while the peak from Lead Pb(4f<sub>7/2</sub>)
(at 138.1 eV) appears at a lower energy than expected from PbI₂ (138.5 eV) and MAPbI₃ (138.7 eV). It may be possible that a small contribution from unbound lead contributes to the Pb(4f<sub>7/2</sub>) peak; we indeed observe some excess intensity on the low energy side of the peak. Note that the measured binding energies are consistent with chemically bonded I and Pb atoms. Let us now look at figure 13c that shows a close-up on the 400 eV region. One clearly identifies the weak Pb(4d) lines, at 412.7 eV (4d<sub>5/2</sub>) and 434.5 eV (4d<sub>3/2</sub>). These energies are moreover consistent with PbI₂. Here, the key observation is the absence of any N(1s) contribution. Finally figure 13d shows the C(1s) energy region and confirms the absence of any C(1s) contribution in the spectrum. Although XPS probes only the top few layers, GIFAD demonstrates a homogeneous structure over the entire film thickness. We can thus conclude that, as for the thinner layers, MAI does not co-adsorb with PbI₂ in the co-evaporation mode.

<figure><img src="image_14.png" /><figcaption>**Figure 11.** Narrow energy high-resolution spectra of major core level contributions from a) Iodine, b) Lead. c) shows the Pb(4d) minor contributions and absence of N(1s) line, d) shows the absence of C(1s) line.</figcaption></figure>
### 3. Discussion

The central question that arises from these observations is why does the perovskite phase fail to form? To match some of the deposition parameters reported in the literature³³,³⁴, we varied the deposition rate from 0.01 ML/min to 0.6 ML/min and the MAI/PbI₂ flux ratio from 0.5 to over 50. Whether MAI and PbI₂ are evaporated simultaneously or sequentially, MAI appears to not adsorb effectively except on the bare silver substrate, as described above. We equally explored various substrate temperatures during both deposition and post-deposition annealing. In all cases, the grown layer consistently shows a hexagonal lattice similar to that of pure PbI₂. Although the silver substrate is expected to strongly influence the overlayer, this effect seems to vanish above 10 ML and should have a negligible influence for a thickness of 50 ML.

In the work of Pellegrino et al.³⁵, MAPbI₃ layers were successfully produced by dipping a thin crystalline PbI₂ film into a solution containing MAI. The resulting film shows XRD diffraction from MAI, although with a shift of the peaks, but not from PbI₂ whose diffraction signal was initially observed. They suggest that MAI intercalates in the van der Waals gap between the PbI₂ monolayers and initiates a chemical reaction that strongly disturbs the PbI₂ network. A subsequent heating at 100 °C for 30 min destroys the MAI arrangement; further heating at 100 °C for 180 min finally produces the MAPbI₃ layer. This transient MAI intercalation inside PbI₂ was reproduced by molecular dynamics simulations³⁶ with a valuable insight into the atomic mechanism. The collapse of the PbI₂ layers results from the nucleophilic attack by iodine migrating from MAI to the PbI₂ monolayers. Subsequent PbIₓ complex formation and rearrangement eventually lead to the formation of MAPbI₃. If such intercalation also occurs in our experiments, then we should probably expect to observe a faster decoupling between the MAI buffer layer and the co-deposited layer, in comparison to deposition of PbI₂ alone. As shown in Figure 8, the substrate effect extends up to the eighth monolayer whether or not we co-evaporate MAI and PbI₂. In all cases, if MAI does intercalate, then the nucleophilic reaction is not taking place.
A comprehensive literature survey on the vacuum deposition of MAPbI$_3$ leads to the conclusion that the ratio of MAI/PbI$_2$ evaporation rates is typically below 10. Assuming a significantly low sticking coefficient of MAI on the co-evaporated layer in our UHV conditions, we have explored MAI/PbI$_2$ flux ratios exceeding 50 without any noticeable effect. The literature survey identifies two distinct methods for the growth of MAPbI$_3$ in vacuum. One, which we refer to as the UHV approach, consists in the growth of ultrathin layers on crystalline metals at vacuum base pressures in the $10^{-10}$ mbar range or lower (e.g., see $^{33,37}$). The other, which we call the high-vacuum (HV) approach, focusses mostly on the growth of thicker layers of MAPbI$_3$, with a thickness of a few hundred nm, in a high vacuum environment (e.g., see$^{38,39}$). With the UHV approach, C. Lang and his colleagues$^{37}$ achieved deposition of MAPbI$_3$ on Au(111) under conditions very similar to our experiment, by co-evaporation of MAI and PbI$_2$ at respective temperatures of 105 °C and 300 °C. Remarkably, they successfully obtained a monolayer of stoichiometric MAPbI$_3$ on a MAI rich buffer layer. The key difference to our methodology is that their substrate is cooled down to 110 K during deposition to favor MAI adsorption and reaction with PbI$_2$. A subsequent annealing is performed at room temperature. Such a processing is not possible in our setup due to our inability to cool down our sample.

At variance, deposition of MAPbI$_3$ in high vacuum, with base pressures spanning from $10^{-7}$ mbar to $10^{-5}$ mbar, is typically accomplished with the substrate held at room temperature. For instance, Wang at al.$^{34}$ deposited a 60 nm thick MAPbI$_3$ layer, with a deposition rate of 1.1 Å/min for PbI$_2$, i.e. in the same order of magnitude than the growth rate of 3.86 Å/min derived from Figure 9. The base pressure in the vacuum chamber is in the $10^{-7}$ mbar range. In these conditions, the water partial pressure amounts to a large fraction of this pressure since the vacuum chamber is not baked. We should recall that for adsorbate-surface systems characterized by a sticking coefficient of unity, a partial pressure of $10^{-6}$ mbar yields a growth rate of 1 ML/s, so much larger than the reported MAPbI$_3$ growth rate. It is then understandable that a non-negligible amount of water covers the growing surface. A complementary study from the same group$^{40}$ demonstrated that water molecules, whose partial pressure is monitored during the deposition process, are actually incorporated into the growing film.
Along the same line, Li et al.⁴¹ achieved a PCE of 20 % for a solar cell with an active MAPbI₃ layer grown by co-deposition in vacuum (base pressure of 8.10⁻⁶ mbar), with a growth rate and a flux ratio comparable to those used in Figure 9. A conclusion that one could draw from these studies, and the two possible routes for achieving MAPbI₃ crystallization by vacuum evaporation, relates to the two pathways for favoring MAI reactivity with PbI₂. In UHV conditions, this is achieved by cooling down the substrate; whereas in HV conditions, water is likely the key ingredient that favors MAI adsorption and possibly reaction with PbI₂.

Relying on this assumption, we take a closer look at the phase transition observed at the start of co-evaporation on the MAI buffer layer (Figure 2). We recall that the relative intensities of the diffraction peaks are, in first approximation, a fingerprint of the electron density distribution within the lattice; any change of this quantity will modify the relative peak intensities. The data is reproduced in Figure 12a. Figure 12b shows quantitatively the time evolution of the absolute intensities within the 0 to 3 diffraction orders. The most spectacular evidence is the ultrafast reorganization of the layer structure triggered by minute amounts of adsorbed PbI₂ molecules. This reaction is particularly visible in the peak intensity of the ±3 orders at t = 90 s. At this point, we brought only 2% of the PbI₂ quantity required to complete the first hexagonal monolayer (at t = 4300 s, see Figure 2). One can also notice the peak intensity of orders ±1 at t = 220 s. in Figure 12b. These variations exclude the hypothesis of a hexagonal overlayer growing on top of the MAI buffer layer, in which case the intensity of all diffraction peak, except the 0th order, would drop monotonously while keeping constant relative values.

Another hint on the nature of the underlying transformation is provided by the evolution of the period measured by GIFAD, displayed in Figure 12c. Alike the 3rd diffraction peak intensity, this period initiates an instantaneous evolution with a sudden drop of 2.3% nearly immediately upon PbI₂ adsorption. While the diffraction signal of the square lattice fades out (Figure 12a), the growing hexagonal phase only shows the specular peak (0th order) since the beam is no more aligned with the crystallographic directions of the new lattice. The diffraction pattern of the hexagonal lattice is observed following a
<figure><img src="image_15.png" /><figcaption>**Figure 12.** a) Time evolution of the diffraction along the Ag [110] direction, same data as in Figure 2. b) Absolute intensities of the 0 to -3 diffraction orders around the square-to-hexagonal transition. c) Period seen by GIFAD as derived from the data in a, for comparison the substrate lattice period along this direction is 2.89 Å.</figcaption></figure>

15° azimuthal rotation of the sample (Figure 2). Complementary information on this square-to-hexagonal transition is gained when the growth is monitored along the substrate Ag[100] direction. The corresponding time evolution of the diffraction pattern is shown in Figure S6a. In this case, one observes that the diffraction signal remains visible throughout the square-to-hexagonal transition, meaning that the long-range order is preserved. Another demonstration against a hexagonal layer growing on top of the MAI buffer layer is provided by the fact that the intermediate spectrum,
measured halfway between the shutter opening and the completion of the hexagonal phase, cannot
be reproduced by a linear combination of the diffraction spectra of the initial and final phases (Figure
S6b). This complete or partial conversion of the MAI square phase is further supported by the near
disappearance of the c(2x2) structure in Figure 5b, whereas the contribution from the Ag substrate is
clearly visible.

From all these observations, we draw two major conclusions: (i) the co-evaporation step does not lead
to the growth of a hexagonal overlayer on top on the square MAI buffer layer but rather transforms
gradually the latter into a Pbl₂ rich hexagonal layer. (ii) The transition is of the displacive type, i.e. the
lattice atoms undergo a coherent motion leading to a change of symmetry. This is in contrast to order-
disorder phase transitions where atoms move in an incoherent manner, producing a total loss of the
diffraction signal. Put together, these two statements demonstrate that the growing layer results from
a rearrangement of the MAI buffer layer triggered by the adsorbed Pbl₂, as sketched in the Figure 13.

<figure><img src="image_16.png" /><figcaption>**Figure 13.** Two possible schemes describing the growth of the hexagonal structure on the MAI buffer layer. The top scenario of a new hexagonal overlayer covering the square buffer layer is incompatible with the GIFAD data of Figure 12. Instead, a displacive phase transition transforms the MAI buffer layer into a Pbl₂ rich layer.</figcaption></figure>
This observation is even more remarkable knowing that the measured signal is averaged over a 0.15 x
8 mm² area, the 8 mm length being limited only by the size of the sample.

The observed high sensitivity of the MAI buffer layer structure to dilute amounts of Pbl₂ molecules is
consistent with the MAI-Pbl₂ pseudo-phase diagram published by Song et al.⁴² Starting with a pure MAI
layer, they followed the transformation induced by the addition of increasing amount of Pbl₂, followed
by annealing up to 190 °C. They also observed that 2% Pbl₂ concentration is sufficient to convert the
crystalline MAI phase to what they call Low Dimensional Perovskites, with crystalline property. Our
finding on the displacive character of the phase transition transforming MAI to Pbl₂ is also consistent
with the degradation dynamics of MAPbl₃ observed in a Transmission Electron Microscope (TEM)²⁵.
Because of the electron irradiation, the MAPbl₃ layer starts to decompose within seconds and gradually
produces pure Pbl₂ while preserving a crystalline order. The nature of the transient phases is identified
by subtle changes in the Selected Area Electron Diffraction (SAED) pattern. These two previous studies
demonstrate that the MAI+Pbl₂ mixture can lead to a crystalline arrangement independently of the
relative concentrations. One can naturally ask whether the intermediate phase, described by the
plateaued value of the lattice parameter shown in Figure 12c, corresponds to MAPbl₃. If the growth is
interrupted near t = 700 s, the structure is indeed stable. The corresponding in-plan primitive cell is
square with a parameter of 4.25 Å, very close to the c(2x2) structure of the MAI buffer layer. These
crystallographic properties however do not match those of MAPbl₃. The reason might be that not
enough iodine is available from the MAI buffer layer because of the strong interaction with Ag, at the
origin of its c(2x2) arrangement. This hypothesis is confirmed by the work of Pellegrino et al.³⁵ who
noticed that the transformation of Pbl₂ layers into MAPbl₃ goes through an intermediate phase that is
only observed for a ratio Pbl₂/MAI of 1/3, corresponding to a I/Pb ratio of 5. This necessary Pbl₂/MAI
ratio is consistent with that used in the single step solution processing followed by annealing at 100°C.
Note that an I-deficient structure, identified as MAPbl₂.₅, and further MA-deficient phases (MAyPbl₂.₅-
x, with 0 ≤ y ≤ 1 and 0 ≤ x ≤ 0.5) have been observed by SAED during MAPbl₃ degradation in a TEM²⁵.
4. Conclusion

We have explored the conditions of MAPbI3 growth by evaporation from two solid sources (MAI and PbI2) in UHV on Ag(001). By the combination of real-time monitoring of the growing layer by GIFAD and complementary analysis by multiple techniques, we demonstrate that the growth of stoichiometric layers of MAPbI3 is not possible in such a clean environment and with the substrate maintained at room temperature. Whether sequential evaporation or co-evaporation is used, the highly volatile MAI does not adsorb on the surface except on the bare Ag substrate. While the substrate effect is visible up to 8 ML of PbI2 on top of the MAI buffer layer, the growth of a much thicker layer by co-evaporation only produces pure PbI2. Nonetheless, we bring evidence of a reaction between the MAI of the buffer layer and subsequently adsorbed PbI2 molecules, with a remarkable and complex phase transition that exhibits a rapid reorganization with only 2% PbI2. This phase transition, displacive in nature, gradually converts the in-plane square lattice of the MAI buffer layer onto an intermediate structure that eventually leads to a PbI2 rich hexagonal layer. These results represent a major challenge for Molecular Dynamics simulations. By showing a very low reactivity between MAI and PbI2, these results suggest that a third chemical compound might be necessary to activate the process. The effectiveness of vacuum evaporation of MAPbI3 in high vacuum environments suggest that water is the key ingredient. As a perspective, the hypothesis mentioned above may be verified by introducing a controlled amount of water during the deposition process. Alternatively, much higher MAI vapor pressures, and correspondingly much higher MAI/PbI2 flux ratios, can be used to enhance the reaction kinetics for producing MAPbI3 layers while keeping a high level of purity. Such a process could effectively be performed and monitored with our new High-Pressure version of GIFAD,43 at working pressures up to 10-2 mbar. From the more fundamental point of view, applying our experimental approach to the low temperature growth process reported for instance by She et al.33 could provide valuable insights into the low-temperature dynamics of the organic species and the expected phase transition during the warming-up phase towards room temperature.
## ASSOCIATED CONTENT

### Supporting Information.

Mass spectrometry measurements; structural properties of the MAI buffer layer; assignment of peaks in the HREELS data; GIFAD reflectivity and layer growth mode; relationship between GIFAD data and lattice properties, comparison between the diffraction patterns measured on the 50 and 11 ML thick layers; growth sequence probed along the Ag[100] direction.

## AUTHOR INFORMATION

Corresponding Authors

*E-mail:
abdallouche89@gmail.com
hocine.khemliche@universite-paris-saclay.fr

## Acknowledgement

The authors gratefully acknowledge the Lebanese Association for Scientific Research (LASER, Tripoli, Lebanon) for their financial support through a scholarship for A. R. Allouche. This work received funding from Investissements d'Avenir, LabEx PALM (ANR-10-LABX-0039-PALM) and from the French National Research Agency (ANR) under the ANR-24-CE30-1014-01 project.

With the aim of publishing it in open access, the author has applied a CC-BY 4.0 open access license to any manuscript accepted for publication (AAM) resulting from this submission.
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