# Intragranular thermal fatigue of Cu thin films: Near-grain boundary hardening, strain localization and voiding

K. Hlushko, T. Ziegelwanger, M. Reisinger, J. Todt, M. Meindlhumer, S.

Beuer, M. Rommel, I. Greving, S. Flenner, J. Kopeček, et al.

► To cite this version:

K. Hlushko, T. Ziegelwanger, M. Reisinger, J. Todt, M. Meindlhumer, et al. Intragranular thermal fatigue of Cu thin films: Near-grain boundary hardening, strain localization and voiding. Acta Materialia, 2023, 253, pp.118961. 10.1016/j.actamat.2023.118961. hal-05080032

HAL Id: hal-05080032

https://hal.science/hal-05080032v1

Submitted on 22 May 2025

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Full length article

# Intragranular thermal fatigue of Cu thin films: Near-grain boundary hardening, strain localization and voiding

<figure><img src="image_3.png" /></figure>

K. Hlushko<sup>a,b</sup>, T. Ziegelwanger<sup>a</sup>, M. Reisinger<sup>c</sup>, J. Todt<sup>a</sup>, M. Meindlhumer<sup>a</sup>, S. Beuer<sup>d</sup>, M. Rommel<sup>d</sup>, I. Greving<sup>e</sup>, S. Flenner<sup>e</sup>, J. Kopeček<sup>f</sup>, J. Keckes<sup>a,*</sup>, C. Detlefs<sup>g</sup>, C. Yildirim<sup>g</sup>

## ARTICLE INFO

Keywords:

Copper

Dark field X-ray microscopy

Residual stress

Thermal fatigue

Thin film

## ABSTRACT

In order to obtain a fundamental understanding of the phenomena accompanying thermomechanical fatigue of Cu metallization used in power electronics, as well as the resulting deterioration of electric properties, there is a need to assess intragranular microstructure and strain evolution within individual Cu grains and near grain boundaries. Here, synchrotron dark field X-ray microscopy (DFXM) is used to characterize as-deposited and 5 × 10⁴ times thermally-cycled 20 µm thick Cu films. The cycling was performed using a dedicated test chip in the range of 100–400 °C applying a heating rate of 10⁶ K/s. The thermomechanical treatment results in severe shear deformation of Cu grains, film surface roughening, gradual grain microstructural refinement, the emergence of microscopic voids preferably at high angle grain boundaries (HAGBs) and finally in the voids’ percolation, as revealed by in-situ and ex-situ scanning electron microscopy. DFXM provides experimental evidence that mosaicity of Cu grains, residual tensile and compressive elastic strain concentrations and full width at half maximum of Cu 111 reflections increase simultaneously in the vicinity of the HAGBs. The latter is interpreted as resulting from vacancy condensation in front of the HAGBs, after dislocations have moved across cycled grains and partly annihilated. Moreover, the observed HAGB decohesion and gradual void formation are correlated with the supposed hardening of regions near HAGBs, which thus lose their ductility during cyclic elasto-plastic deformation. The results are complemented with the ex-situ X-ray nanotomography data, which document the voids’ percolation across the film’s thickness with a “crack” width of up to ~2 µm. In general, the study identifies the inhomogeneous intragranular microstructural refinement and a gradual condensation of structural defects near HAGBs as driving forces for void formation in thick Cu metallizations during fast thermo-mechanical fatigue.

## 1. Introduction

The main challenges of microelectronic technology comprise the miniaturization of electronic components coupled with constantly increasing demands on their performance, like higher switching frequencies and power densities [1,2]. These conflicting requirements stimulate the development of novel device designs, fabrication processes and optimized material choice. Since microelectronic devices are composed of very dissimilar materials, such as semiconductors, metals,

ceramics and plastics, the increase of power densities in ever more miniaturized components results in highly localized temperature hot-spots and consequently in the generation of cyclic micro- and nano-scopic thermomechanical loads, as a result of the mismatch of coefficients of thermal expansion (CTEs), often in a complex interplay of multiple adjoining components [3,4].

Because of its excellent electric and thermal properties, Cu metallization is commonly used in microelectronics for charge transfer and as a heat sink [5]. Thermal cycling, however, causes fatigue phenomena,

<footnote><sup>a</sup> Department of Materials Science, Montanuniversität Leoben, Leoben, Austria</footnote>

<footnote><sup>b</sup> Materials Center Leoben Forschung GmbH, Leoben, Austria</footnote>

<footnote><sup>c</sup> KAI – Kompetenzzentrum Automobil- und Industrieelektronik GmbH, Villach, Austria</footnote>

<footnote><sup>d</sup> Fraunhofer Institute for Integrated Systems and Device Technology IISB, Erlangen, Germany</footnote>

<footnote><sup>e</sup> Institute of Materials Physics, Helmholtz-Zentrum Hereon, Geesthacht, Germany</footnote>

<footnote><sup>f</sup> FZU - Institute of Physics, Czech Academy of Science, Praha, Czech Republic</footnote>

<footnote><sup>g</sup> European Synchrotron Radiation Facility, Grenoble, France</footnote>

<footnote>* Corresponding author.
E-mail address: keckes@unileoben.ac.at (J. Keckes).</footnote>
which can result in thermomechanical damage of the metallization, localized heat dissipation, thermal runaway, deterioration of electric properties and finally in device dysfunction [6]. The thermomechanical behavior of Cu thin films and its correlation to film thickness, microstructure, residual stresses, elastic anisotropy, surface roughening, grain boundary (GB) voiding, electrical properties, electro-migration and diffusion into the substrate have been extensively studied for decades [6–12].

The thermal fatigue of Cu is strongly film thickness-dependent, especially as a consequence of the strong elasto-plastic anisotropy of Cu and the gradual evolution of crystallographic texture [13,14], which evolves progressively across the film thickness from <111> to <100>-type fiber texture, with a transition at a thickness of $h_f \approx 1$ µm [7,15]. Typically, Cu films with $h_f$ in the range of 0.1–5 µm exhibit thermomechanical damage in terms of localized extrusion/intrusion events and severe surface roughening [12,16,17].

To *improve the reliability* of Cu metallization in power electronics, a relatively recent approach considers an increase of $h_f$ up to ~20 µm, which results in an increase of heat capacity, in a longer operational lifetime and even in a higher maximum sustainable peak temperature [16]. The thermomechanical fatigue of such thick Cu films has been studied using an infrared furnace, a fast laser cycling system and dedicated testing chips. Previous studies indicated that the fatigue behavior depends decisively on the applied heating rates (16,920 vs. $10^6$ K/s) [18–21], the applied thermal pulse base and maximal temperatures (~145 vs. 445 °C) [22,21] and the number of applied cycles [16,20,22, 23]. Bigl et al. [16] have shown that, by increasing $h_f$ from 5 to 20 µm, the average grain size and the fraction of (100) texture component increase, while the mean in-plane elastic moduli and the fractional twin boundary length decrease after ~$10^3$ applied cycles. These effects are accompanied by a continuous enlargement of (100) oriented grains and simultaneous shrinkage of differently oriented grains, which gradually reorient and deform by severe shear on multiple glide planes [16]. The fatigue processes are usually accompanied by a continuous accumulation of surface damage, as a result of glide step formation in the vicinity of GBs and hillock-like structures at the GBs, whose lateral surface roughness appears to scale with $1/h_f$ [16,24]. Kleinbichler et al. [22] reported the formation of microscopic voids at GBs and triple points, whose concentration and volume increased with the increasing number of thermal pulses. Similarly, Moser *et al.* [21] applied up to $1.2 \times 10^4$ cycles and observed an increase in electrical resistance due to gradual formation of intergranular cracks, which was correlated to an increase in heat pulse base-temperate and a decrease of pulse length. Remarkably, for short heat pulses of ~0.2 ms, intergranular cracking was reported, whereas for longer heat pulses, a change in the dominant degradation mechanism from cracking to more pronounced dislocation-driven plasticity was observed [21].

In general, previous studies on the thermal fatigue of Cu metallization indicated complex changes to the microstructure and surface roughness of Cu films, triggered mainly by cyclic tensile-compressive thermal stresses and surface diffusion, and controlled by the initial film thickness, microstructure, maximal applied temperature, heating pulse rate and base temperature and the number of applied cycles [6,8, 12,14,16,17,21,22,25]. The main drawback of practically all previous studies, however, resides in their spatial resolution. The common experimental techniques provide either (i) only volume-average data, like on the stress-temperature behavior determined by in-situ wafer bow and/or X-ray diffraction, or (ii) surface characteristics of the films, like scanning electron microscopy (SEM) and electron backscatter diffraction (EBSD) analyses of the films' surfaces and cross-sections. Moreover, since EBSD is very sensitive to surface quality and requires a low roughness, the approach is applicable primarily to structures with a limited number of applied fatigue cycles, up to ~$10^3$ [16,19,20].

In order to obtain a fundamental understanding of the micro- and nanoscale degradation processes accompanying thermal fatigue of Cu thin films, there is a need to take a further step and (i) to assess residual

strain and microstructural evolution within individual Cu grains, (ii) to understand the origins of changes in GB morphology and accompanying voiding phenomena and (iii) to analyze films after a higher number of applied thermal cycles. The main motivation for this work is to experimentally study intragranular microstructure across individual Cu grains after thermal cycling and to assess lateral distributions of strains, crystallite orientation and defect density. An ex-situ characterization was performed mainly using dark field X-ray microscopy (DFXM) on as-deposited and $5 \times 10^4$ times thermally cycled Cu grains. The results are complemented with data from X-ray nanotomography, EBSD and in-situ SEM. The main experimental findings indicate (i) a microstructural refinement of Cu grains, which is more pronounced near GBs, and (ii) an increase in residual elastic strain and defect density near GBs. Therefore, the observed film voiding is correlated with a supposed embrittlement of near-GB regions, promoting gradual decohesion at GBs. Additionally, this first DFXM study of thermal fatigue in metallic films also demonstrates the potential of the relatively novel X-ray microscopy technique, to assess relevant intragranular phenomena in thin metallizations, especially the intragranular (2nd order) strains and their correlation to void formation and other microstructural damage.

## 2. Experiment

The aim was to experimentally study the thermomechanical behavior of as-deposited and $5 \times 10^4$ times cycled Cu metallization layers. For this purpose, dedicated test chips were used, as described in detail in the work of Moser et al. [18]. The chips were fabricated using an industrial-scale semiconductor manufacturing process on Si (100) wafers with a diameter of 300 mm. The chips' multilayer stack consists of a 120 µm thick Si substrate, a built-in polycrystalline Si heating plate, an Al-based temperature sensor as well as a 20 µm thick Cu metallization layer, electrochemically-deposited on a previously sputtered 300 nm thin WTi interlayer. In order to stabilize the microstructure of the metallization layer, the entire test chips were annealed at 400 °C in forming gas atmosphere for 30 min. The electrochemical-deposited Cu metallization studied in the present work was processed using an electrolyte, which results into an almost impurity free Cu film with an average grain size of about 20 µm equivalent circle diameter after the anneal. Fig. 1a shows the design of the test chip with visible Cu metallization pads. Only the large central pad with dimensions of $700 \times 700 \times 20$ µm$^3$ is heated during experiments, while the small pads on the sides serve as electrical contacts for the functional layers described above.

Thermal cycling was performed in forming gas atmosphere within a temperature range of 100–400 °C, using a heating rate of $10^6$ K/s, a heating pulse length of 200 µs and a repetition frequency of 1 Hz, as described in Ref. [18]. Considering the CTEs of Cu and Si of 17 and 2.7 $10^{-6}$ K$^{-1}$, respectively, the heating pulse amplitude and rate correspond to a temperature-induced strain amplitude and strain rate of ~0.4% and ~0.8 s$^{-1}$, respectively [21].

In Fig. 2a, one exemplary measured temperature pulse applied to the test chip is presented. The fast thermo-mechanical pulsing experiments with $5 \times 10^4$ cycles were performed on a total number of 5 chips, all of which indicated practically the same degradation behavior when inspected in SEM. For further DFXM analysis, two representative chips cycled $5 \times 10^4$ times (Figs. 1b,6,8,9 and 10) and one pristine chip (Figs. 4,5) were selected. Additionally, one chip was cycled only $10^4$ times and further analyzed by EBSD (Fig. 7) in order to assess early stages of the thermomechanical damage.

A complementary thermomechanical characterization of a representative (planar) Cu thin film deposited on Si(100) substrate was performed using a wafer curvature system in the temperature range of 25–400 °C applying a heating and cooling rate of 5 K/min (Fig. 2b). For this purpose, a kSA multiple optical beam sensor (MOS) (k-Space Associates, Inc., Dexter, Michigan) device was used to quantify the stress dependence on the (relatively slowly) changing temperature.
<figure><img src="image_4.png" /><figcaption>**Fig. 1.** (a) An optical micrograph of the test chip used to thermally cycle the rectangular Cu pad of 700×700×20 µm³ in its center [18]. In (b) an exemplary SEM micrograph of the surface of the Cu pad after 5 × 10⁴ applied cycles is presented. The arrow indicates a grain which was further analyzed using DFXM and X-ray nanotomography (cf. Figs. 8 and 10).</figcaption></figure>

SEM and EBSD characterization of surfaces of the as-deposited and cycled films was performed using a LEO 1525 (Zeiss, Inc., Oberkochen, Germany) scanning electron microscope, applying an acceleration voltage of 20 kV and a 120 µm aperture size. The EBSD data were collected only from the as-deposited film areas of 175×130 and 75×50 µm² using steps of 0.75 or 0.28 µm, respectively. A cross-section of a Cu thin film cycled 10⁴ times (Fig. 7) was fabricated using a Zeiss CrossBeam 1540XB (Zeiss, Inc., Oberkochen, Germany) focused ion beam (FIB) workstation and also investigated by EBSD. EBSD on the cycled Cu thin film was performed with a Tescan MAGNA (Tescan Orsay, Brno, Czech Republic) SEM operated at 30 kV accelerating voltage and 6 nA beam current equipped with an e⁻-Flash<sup>FS</sup> EBSD-detector (Bruker, Billerica MA, USA). For acquisition, a step size of 75 nm was used. The post-processing of the EBSD data the Bruker Esprit software version 2.3 was used, where the minimum grain size was set to 10 px, while a grain tolerance angle of 15 deg was applied. Missing pixels within a continuous matrix were assumed to have an orientation similar to the matrix' orientation.

The lateral strain distribution and mosaicity within the as-deposited and 5 × 10⁴ times cycled Cu thin films were characterized at beamline ID06 of the European Synchrotron Radiation Facility (ESRF) using DFXM [26]. For the experiment, an X-ray beam was used with a photon energy of 17 keV selected by a Si(111) double crystal monochromator. The incident beam was focused by means of two-dimensional (2D) transfocators with Be compound refractive lenses (CRLs) having 14 lenslets, resulting in a 200×200 µm² (horizontal × vertical) spot size on the samples, which was substantially larger than the diffracting Cu grains of interest. Diffraction intensities of Cu 111 reflections (Fig. 3) collected from (randomly selected) Cu grains (within as-deposited and 5 × 10⁴ times cycled films) were magnified by an X-ray objective comprising 88 Be lenses. The projections of the grains were captured by a 2D indirect detector with an effective pixel size of 0.75 µm (Fig. 3). The magnification of the X-ray lens of ~18.1 provided an effective pixel size at the sample of ~42 nm and a resolution of ~100 nm [27,28]. The recorded images are projections integrated along the beam path through the diffracting Cu grains and therefore represent grain thickness-averaged microstructural and strain information (cf. Fig. 3).

<figure><img src="image_5.png" /><figcaption>**Fig. 2.** (a) A time profile of an exemplary power pulse with a length of 200 µs applied to the Cu pad on a chip (Fig. 1a) and the resulting temperate pulse (in the range of 100 – 400 °C, with a heating rate of 10⁶ °K/s) measured by a temperature sensor. (b) A complementary stress-temperature dependence, measured on a *compact Cu thin film* using a wafer curvature system (using a heating and cooling rate of 5 K/min), shows the stress evolution during the first and the second thermal cycle.</figcaption></figure>

Two representative thermally-cycled Cu thin films and one pristine (not cycled, as-deposited) Cu film on three chips (Fig. 1a) were used for the DFXM analysis. On the heated chips, several grains were analyzed by DFXM. Hereafter, DFXM results from only two representative Cu grains (each from a separate chip) are presented (Figs. 8,9). Similarly, several Cu grains within the pristine film were analyzed by DFXM and results from one representative grain are presented in Fig. 5.

The *crystallographic misorientation* (intragranular grain mosaicity) was assessed by scanning one (ω) or two (ω, χ) sample rotation axes while the scattering angle 2θ was fixed to the Cu 111 reflection position. The orientation scans were recorded with a step size of 0.02° over a maximum range of up to 1°, depending on the grain. For *strain scans*, the sample angle ω, the scattering angle 2θ (by collective motion of detector and the objective) were scanned, to capture a ω-2θ mesh, at a constant χ angle (cf. Fig. 3). The results from the strain scans (ω-2θ) show intragranular variations of the lattice spacing of Cu {111} crystallographic planes, whereas the orientation (ω-χ) scans at a constant 2θ angle indicate the spatial variation in the {111} plane normals' orientations. The latter may be caused by either tilt or twist of the crystallographic lattice or due to the presence of shear strains.

For data evaluation, we considered the intensity at a pixel coordinate as function of the goniometer angles across the scan. The center-of-mass positions then form 2D maps of strain and crystallite orientation. The maps of integrated intensity provide reference images of a grain's morphology. Details of the generation of these maps can be found in Ref. [26].

Since the Cu film cycled 5 × 10⁴ times exhibited micro-cracks at the surface (cf. Fig. 1b), X-ray nanotomography was carried out to analyze the three-dimensional (3D) volumetric morphology in the vicinity of the specific grain indicated by the arrow in Fig. 1b. For this purpose, a
<figure><img src="image_6.png" /><figcaption>**Fig. 3.** Schematic diagram of the DFXM experimental set-up [29]. The diffracted signal from a randomly selected Cu grain within the rectangular Cu pad (Fig. 1a) is magnified through an X-ray objective CRL located 275 mm behind the sample and then collected by a 2D detector ~5 m downstream from the sample. ω-χ and ω-2θ mesh scans were used to tilt the sample and the sample-CRL-CCD pair in order to evaluate 2D distributions of mosaicity and strain, respectively, which were averaged along the beam path in the grain. The grain's reciprocal space vectors Q₁₁₁ were oriented approximately ~25° out of the Cu pad's in-plane direction.</figcaption></figure>

cylindrical pillar with a diameter of ~50 µm and a height of ~35 µm containing the grain was prepared. The pillar preparation was carried out in a Helios G4 PFIB CXe system from ThermoFisher Scientific. A carbon/platinum protective layer was first deposited onto the Cu film's surface. Subsequently, the circular milling of the pillar was carried out by means of Xe ion beam sputtering at 30 kV and high ion beam current. After lift-out, the pillar was fixed to a sample carrier pin. Finally, the pillar was polished by using low ion beam currents. The sample preparation procedure is presented in Supplementary Fig. 1.

Afterwards, the pillar was mounted on a steel tip and then analyzed at the imaging beamline P05 of the storage ring PETRA III at DESY in Hamburg, operated by Helmholtz Zentrum Hereon. The pillar was imaged using an X-ray energy of 8.8 keV in absorption mode, applying a full field transmission X-ray microscopy (TXM) setup [30]. The TXM was equipped with a Fresnel Zone plate (FZP) of 150 µm in diameter possessing an outermost zone width of 50 nm. In combination with beam shaping optics and field sizes of 50×50 µm², an effective pixel size of 14.8 nm was achieved. Thanks to the unique geometry of the instrument, a Hamamatsu X-ray detector with 6.5 µm physical pixel size was placed at a sample to detector distance of 20.5 m allowing for a sufficient magnification. The collected data were reconstructed using the TomoPy reconstruction tools [31] and finally a pixel binning of 2 was applied, resulting in an effective voxel size of ~26 nm³.

### 3. Results and discussion

#### 3.1. As-deposited Cu thin film

The initial microstructure of the as-deposited 20 µm thick Cu film was analyzed using DFXM and afterwards by SEM and EBSD. EBSD data in Fig. 4b and in Supplementary Figs. 2 and 3 indicate an average in-plane Cu film grain size in the micrometer range and a (111) preferred orientation of the crystallites separated by both low and high angle grain boundaries (LAGBs and HAGBs). This is documented also by pole figures calculated from the orientation data of Fig. 4b, which are presented in Supplementary Fig. 2. In Supplementary Fig. 3, several representative SEM micrographs from a cross-section of the as-deposited Cu thin film indicate the absence of voids. Additionally, the EBSD data from the same film cross-section in Supplementary Fig. 4 document a relatively low concentration of LAGBs within the Cu grains.

In the center of the orientation maps in Fig. 4b,c, there is a large (100) oriented area with twins oriented approximately (111) in out-of-plane direction. Fig. 4d shows the in-plane orientation of the grains in this area, which are separated by LAGBs with grain misorientations (15°. Since this in-plane orientation of the (111) and (100) out-of-plane oriented grains appears very similar (Fig. 4d), DFXM (Fig. 3) was used to analyze exactly this film region by using a relatively small angular range for (ω, χ) and (ω - 2θ) scans. As already mentioned in Section 2, the

<figure><img src="image_7.png" /><figcaption>**Fig. 4.** A SEM micrograph from an exemplary Cu grain within an as-deposited thin film (a) and the corresponding EBSD data showing the out-of-plane (b,c), horizontal in-plane (d) orientation of the crystallites with respect to the film surface. KAM is presented in (e). The scale bars equate to 10 µm. Complementary EBSD data are presented in Supplementary Fig. 2.</figcaption></figure>
diffraction vector was oriented approximately ~25° out of the Cu pad's surface.

The observed preferred appearance of 〈111〉 and 〈100〉 oriented grains within the Cu thin film (Fig. 4 and Supplementary Fig. 2) originates from surface/interface and elastic energy minimization during the film's growth, respectively [15,32]. This behavior is typical for the growth of metal films with a high mobility of adatoms and is denoted as Volmer-Weber growth [32,33]. It starts with an independent nucleation of islands, later sees their coalescence and finally a compact film is formed. In the present case, the EBSD data indicate that the electrochemical deposition of the Cu layer obeyed this type of growth mode. The appearance of twins in Fig. 4 can be attributed to the relatively low stacking fault energy of Cu, which typically results in a high number of twin boundaries [34].

The EBSD data collected from a relatively large film area (Fig. 4b) additionally document that, at the macroscopic scale, the film's mechanical properties are in-plane isotropic and therefore the film will be equibiaxially stressed (by 1st order strain) when heated homogeneously. At the microscale (2nd order strain), however, the differently oriented Cu grains are expected to possess different in-plane residual stresses, which vary abruptly in the film's in-plane direction, due to copper's elastic anisotropy with a Zener ratio of ~3.21.

DFXM results from the thin film area with the preferred 〈100〉 out-of-plane orientation in the center of the SEM and EBSD micrographs (cf. Fig. 4) are presented in Fig. 5. The DFXM 2D distributions express Cu properties averaged along the beam path in the diffracting grain, where the scattering vector $Q_{111}$ directions was oriented ~25° with respect to the film surface. The 2D orientation (ω-χ) maps in Fig. 5a indicate that only a part of the diffracting area (Fig. 4a) was covered by the scanning and that there are many orientation-related artefacts, such as the mentioned tilt of 25°, which makes the grain look differently shaped compared to the EBSD data shown in Fig. 4. One can recognize a

triangular segment indicated by the arrows in Figs. 4a and 5a, as well as the presence of some twins with straight boundaries. The differences between the DFXM and EBSD orientation maps can further be explained by the fact that the SEM and EBSD data (Fig. 4) represent information from only the surface, since the penetration depth of electrons is only a few nm, while the DFXM data (Fig. 5) represent diffraction data averaged along the X-ray beam path, i.e. almost the entire film thickness, since the X-ray absorption length is approx. 20 µm at 17 keV.

In Fig. 5c, a normalized axial X-ray elastic strain (Δ2θ/2θ₀) distribution obtained from a ω-2θ scan shows the intragranular relative variation in the Cu 111 peak's Bragg angle position, with pronounced negatively and positively strained regions, which can be interpreted as respective compressive and tensile elastic strain concentrations within the grain. The strain data show that (i) there is a gradual strain change from compressive to tensile across the grain and, interestingly (ii) both tensile and compressive strain concentrations are localized especially at the interfaces between the twins oriented vertically in Fig. 5c. Again, not the whole grain was covered by this particular ω-2θ scan and regions with possibly even larger strains are thus not depicted in Fig. 5c.

In Fig. 5d, full widths at half maximum (FWHM) of Cu 111 peaks collected during ω-2θ scanning are presented. The results indicate a FWHM increase at grain and twin boundaries, which matches the increased X-ray elastic strain concentrations in Fig. 5c. Generally, the FWHM is sensitive to the size of coherently diffracting domains and also to crystallographic defects like dislocations. In the present case, the simultaneously increased X-ray elastic strains (Fig. 5c) and FWHMs (Fig. 5d) are expected to correlate with a higher concentration of (geometrically necessary) dislocations in these areas. The crystallite size influence on the Cu 111 peak broadening is not supposed to contribute significantly to FWHM increase, especially due to the large size of Cu crystallites along the $Q_{111}$.

<figure><img src="image_8.png" /><figcaption>**Fig. 5.** The reconstructed DFXM data from the large Cu grain in the as-deposited film (cf. Fig. 4a). (a) Results from a mosaicity ω-χ mesh scan, complemented by (b) a kernel-average misorientation map indicate an abrupt variation in the diffraction vector $Q_{111}$ orientation at sub-grains and twin boundaries (in degrees). (c) Relative axial Δ2θ/2θ₀ distributions obtained from ω-2θ scans show the 2nd order residual X-ray elastic strain distribution, which changes continually across the grain with local maxima at the twin boundaries and some HAGBs. The increased strain concentrations correlate well with (d) the higher FWHM values (in degrees), especially at twin boundaries.</figcaption></figure>
## 3.2. Thermally cycled Cu thin films

The thermal cycling of the Cu pad in the test chips (Fig. 1a) resulted in repeated formation of alternating in-plane compressive and tensile thermal stresses as a result of the mismatch between CTEs of Cu and Si. Fig. 2b shows an exemplary stress-temperature curve measured using a wafer curvature system on a planar Cu thin film, applying heating and cooling rates of 5 K/min and inducing maximal tensile and compressive stresses of 100 and –155 MPa, respectively. After the first cycle, which results in some degree of microstructure stabilization, the stress-temperature hysteresis remains relatively unchanged thereafter. Upon heating up and cooling down, the Cu film is first deformed elastically, as indicated by a linear stress-temperature dependence, and then yields in compression during heating and then in tension during cooling, as indicated by the decrease of stress-temperature slopes (Fig. 2b). The decrease of compressive stress in the Cu film at temperatures above ~200 °C (Fig. 2b) can be explained by enhanced dislocation mobility (softening) and a creep mechanism, like a diffusional creep and power law breakdown [35].

In the case of the fast heating experiments on the test chips (Fig. 1a), cycling with a heating rate of $10^6$ K/s is expected to induce similar levels of alternating compressive and tensile stress in the Cu pad (Fig. 2b), which also result in cyclic elasto-plastic deformation of Cu, because the yield stress is also reached repeatedly both in compression and tension during heating up and cooling down, respectively. Moreover, it is supposed that fast thermal heating (Fig. 2a) resulted in the build-up of macroscopically homogeneous equibiaxial thermal stresses within the heated film, also because of the missing epitaxial relationship between the film and the substrate, similar as in the film tested in the wafer-curvature system (Fig. 2b). Currently, however, there is no experimental technique available, which would allow for determining the evolution of thermal stresses and their spatial distribution at a heating rate of $10^6$ K/s and thereby verify this assumption.

Fig. 6 and Supplementary Video 1 show exemplary gradual changes in film surface morphology during 5 × $10^4$ thermal cycles in the temperature range of 100–400 °C, as imaged with a SEM. One can observe an increasing surface roughening of the Cu thin film (Fig. 6b), the formation of glide bands (Figs. 6b,c), severe shear deformation of the grains with extrusion and intrusion phenomena, the formation of valleys at HAGBs (Figs. 6b-d) and, finally, a decohesion at HAGBs resulting from

the coalescence of voids (Figs. 6d-f). A triangularly shaped region to the right of the central area (Fig. 6c) shows an accumulation and arrest of multiple shear bands [14,16,19,36,37], where the grain interior's surface appears unaffected by the thermomechanical treatment up to ~$10^4$ cycles. The three-fold symmetry of the progressing plastic deformation suggests a (111) out-of-plane orientation of the respective grain, which is partly preserved during the thermal cycling. Moreover, the Supplementary Video 1 shows a gradual migration of the HAGBs on the film surface, which is expected to promote the formation of micro-cracks and finally voids in the film. These exemplary data document the mechano-structural complexity of severe thermal fatigue in Cu films. From the SEM micrographs in Fig. 6 it is however not possible to conclude (i) where the voids initiate and (ii) what is the driving force for their formation. Therefore, further cross-sectional EBSD and DFXM analyses were performed (cf. Figs. 7-9).

In order to analyze the early stages of the void formation (Fig. 6d), an ex-situ EBSD analysis was performed on another Cu thin film after $10^4$ thermal cycles, whose cross-section was prepared by FIB milling. Representative SEM and EBSD data document the presence of a near-interface void (a full arrow in Fig. 7a) and a microstructural fragmentation of grains resulting in a relatively high density of LAGBs (Figs. 7b, d). The near-interface void is positioned at one end of a twin structure, which boundaries are indicated by dotted arrows in Figs. 7c,b. It should be noted that the similar high concentration of LAGBs was observed across the entire cross-section of the cycled Cu film. For a comparison, EBSD from an as-deposited Cu thin film cross-section in Supplementary Fig. 4 show no significant occurrence of LAGBs and voids. The ex-situ EBSD results from Fig. 7 document that the occurrence of voids in thermally cycled Cu thin films (Fig. 6) can be correlated also with the presence of twins and the void growth can originate also at Cu film surfaces. Previous studies of Kleinbichler et al. [22] documented the void formation in a Cu film after 1000 thermal cycles preferably at HAGBs and triple points.

The ex-situ SEM micrographs from the Cu film surfaces in Figs. 1b, 6, 8e and 9e indicate severe surface roughening and formation of voids after the 5 × $10^4$ applied cycles. Therefore, it was unfortunately not possible to perform an EBSD on these structures and to evaluate the microstructural changes accompanying the thermomechanical damage. Also in the literature, up to now, EBSD analysis has been reported only for Cu thin film surfaces after a maximum of approx. $10^3$ thermal cycles [16,20].

<figure><img src="image_9.png" /><figcaption>**Fig. 6.** SEM micrographs show the surface of a Cu thin film in (a) as-deposited state and after (c-f) 1k, 3k, 10k, 30k and 50k thermal cycles in the temperature range of 100–400 °C (cf. also supplementary Video 1). The arrows in (c) indicate glide steps within a (111) out-of-plane oriented grain [16,36,37], which remain relatively stable up to ~$10^4$ cycles, whereas other areas show valley formation, voiding and decohesion at HAGBs. The scale bar (a) equates to 10 µm.</figcaption></figure>
<figure><img src="image_10.png" /><figcaption>**Fig. 7.** EBSD analysis of a Cu thin film cross-section after $10^4$ thermal cycles. SEM micrograph of the FIB-polished cross-section (a), the image quality map (b), the out-of-plane orientation map (c) and the kernel average misorientation $\theta_{KAM}$ map (d) show the presence of a near-interface void (indicated by a full arrow) and a relative high density of LAGBs. The dotted arrows in (b) and (c) indicate positions of twin boundaries. The scale bar (a) equates to 10 µm. The EBSD map color key from Fig. 4 applies.</figcaption></figure>

DFXM can circumvent this issue and was therefore used to analyze intragranular thermomechanical damage of the cycled Cu grains shown in Figs. 8e and 9e. The results from 2D orientation scans ($\omega-\chi$) document a variation in lattice plane orientations across the grains' interiors (Figs. 8a and 9a), as indicated especially by the kernel-average misorientation maps in Figs. 8b and 9b. The observed increased mosaicity is in agreement with the EBSD data presented in Fig. 7 and confirms the microstructural refinement of the Cu grains. The main qualitative difference, compared to the DFXM results from the as-deposited grain (Figs. 5a,b), resides however in the micrographs in Figs. 8b and 9b, which document a higher degree of microstructural refinement near GBs.

The axial X-ray elastic strain ($\Delta2\theta/2\theta_0$) distributions obtained from $\omega-2\theta$ scans in Figs. 8c and 9c show relatively negligible X-ray elastic

strains within the Cu grains' interior after $5 \times 10^4$ cycles. Comparably large concentrations and singularities of both compressive and tensile X-ray elastic strains were however observed near the GBs, fanning out towards the interiors in both grains. The increased concentrations of strains of the 2nd order, which alternate along the grains' circumference, may originate from orientation-dependent elastic incompatibility between the particular grain and its neighbors.

Similarly, DFXM analysis indicated a relatively small FWHM of the Cu 111 peak within the two cycled grains' interiors and increased FWHM values close to their GBs (cf. Figs. 8d and 9d). The higher FWHM values (Figs. 8d and 9d) suggest a higher concentration of structural defects in these particular regions of the grains, which can be interpreted as an increased density of nanoscopic structural defects like dislocations. In the

<figure><img src="image_11.png" /><figcaption>**Fig. 8.** The reconstructed DFXM data from a Cu grain cycled $5 \times 10^4$ times (cf. Fig. 1b). (a) Results from a mosaicity $\omega-\chi$ mesh scan complemented by (b) a kernel-average misorientation map indicate a diffuse variation of the diffraction vector $Q_{111}$ orientation across the grain (in degrees), with concentrations at the GBs. (c) Relative axial strain $\Delta2\theta/2\theta_0$ distributions obtained from $\theta-2\theta$ scans show the 2nd order residual X-ray elastic strain distribution across the grain. In (d) FWHM values (in degrees) from the $\theta-2\theta$ scans are presented. In (e), a SEM micrograph of the corresponding grain's surface is shown (cf. also Figs. 1b, 10 and Supplementary Fig. 1).</figcaption></figure>
<figure><img src="image_12.png" /><figcaption>**Fig. 9.** The reconstructed DFXM data from another Cu grain cycled 5 × 10⁴ times. (a) Results from a mosaicity ω-χ mesh scan complemented by (b) a kernel-average misorientation map indicate a diffuse variation of the diffraction vector Q₁₁₁ orientation across the grain (in degree), with concentrations at the GBs. Relative (c) axial strain Δ2θ/2θ₀ distributions obtained from 0-2θ scans show the 2nd order residual X-ray elastic strain distribution across the grain. In (d) FWHM data (in degrees) from the 0-2θ scans are presented. In (e), a SEM micrograph of the corresponding grain's surface is shown.</figcaption></figure>

case of soft Cu thin films, the dislocation movement across grains is dimensionally constrained by HAGBs and the dislocations typically pile up at interfaces [17,24,36,38]. Therefore, the increased FWHM values in Figs. 8d and 9d can be interpreted as an increased dislocation density near HAGBs, as a result of accumulated thermomechanical damage.

The simultaneous increase in crystal mosaicity, X-ray elastic strains and FWHMs in the vicinity of HAGBs (Figs. 8 and 9) documents that these play an important role in thermal fatigue of the observed grains. The inhomogeneous microstructural refinement and the accumulation of elastic strains and structural defects near GBs indirectly indicate also an inhomogeneous hardening of the grains, which is localized mainly near HAGBs, while the grains' interiors supposedly remain comparatively strain-free and soft [14,17,36,39].

The hypothesis of the inhomogeneous hardening across grains is supported also by the observed absolute values and ranges of residual strains and FWHMs. In the as-deposited grain, the 2nd order X-ray elastic strain and FWHM ranges are ~5 × 10⁻⁴ and ~4 × 10⁻⁴ degrees (Figs. 5c and 5d), whereas in the 5 × 10⁴ times cycled grains the values reach up to ~12 × 10⁻⁴ and ~2 × 10⁻³ degrees (Figs. 8c,9c and 8e,9e). In other words, both the 2nd order X-ray elastic strains and the structural defect densities are significantly higher in the thermally cycled films, but accumulated mainly near the HAGBs (Figs. 8 and 9).

### 3.3. X-ray nanotomography on Cu thin film cycled 5×10⁴ times

Since the SEM analysis of the Cu thin film cycled 5 × 10⁴ times revealed large coalesced voids on the films' surface, the large central Cu grain and its vicinity (cf. Figs. 1b and 8e) was isolated using FIB milling and subsequently analyzed using absorption contrast-based X-ray nanotomography. The motivation was to assess the in-plane and cross-sectional morphology of the voids and also to estimate their degree of percolation. The representative data in Fig. 10 show the morphology of a representative void, which spans across the whole film thickness of ~20 µm. The maximal opening of the void is ~2 µm. Supplementary Video 2 shows a sweep through out-of-plane (Fig. 10a) virtual cross-sections, again documenting the high degree of void percolation. Supplementary Video 3 shows the same sweep for the as-deposited Cu thin film, demonstrating the absence of voids in the as-deposited thin film.

## 4. Mechanism of intragranular thermomechanical fatigue of Cu thin films

DFXM provided 2D data on the mosaicity, 2nd order X-ray elastic strain and FWHM distributions averaged across the thicknesses of one as-deposited (Fig. 5), and two Cu grains cycled 5 × 10⁴ times thermally

<figure><img src="image_13.png" /><figcaption>**Fig. 10.** Results from absorption contrast-based X-ray nanotomography on the Cu film cycled 5 × 10⁴ times, showing the voids with a thickness of up to ~2 µm penetrating the whole film thickness (cf. also Supplementary Fig. 1 for the sample preparation and Supplementary Fig. 5). The transparent material is the Cu of the grains separated by the voids shown in blue.</figcaption></figure>

(Figs. 8,9). Together with the SEM and EBSD data from the films surfaces (Figs. 1b,4,6 and Supplementary Video 1) and cross-sections (Fig. 7, Supplementary Figs. 3,4), the reported results allow us to draw valuable conclusions on the origins and mechanisms of the thermomechanical degradation within the 20 µm thick Cu films after 5 × 10⁴ thermal cycles, especially elucidating the reason for the observed HAGB decohesion and void formation (cf. Fig. 6 and Supplementary Video 1).

Previous studies of Cu thin films with a thickness of $h_f > 5$ µm indicated a strong dependence of thermomechanical damage on the number of applied cycles, heating rate and applied temperatures [16,20, 21], while two types of degradation mechanism have been distinguished. Most studies focused on the analysis of Cu films cycled “slowly” using an infrared furnace or a laser, where thermomechanical damage is controlled by dislocation-driven plastic deformation and constrained diffusional creep, which result in roughening of the film surface through formation of extrusions and/or gradual enlargement of (100) oriented grains [16]. When short thermal pulses of 200 µs and fast heating rates of 10⁶ K/s are applied, as in the present case, a dominant cracking of the
Cu films was observed by SEM [21].

The present data complement and extend the knowledge found in literature by (i) increasing the number of applied cycles to 5 × 10⁴ and (ii) by applying DFXM to assess intragranular structural information from the individual Cu grains after the decohesion of HAGBs. First of all, the applied thermal cycling results in microstructural refinement of grain interiors and in the formation of small crystalline domains separated by LAGBs (cf. Fig. 7c-d), which were not observed in the as-deposited grains (cf. Supplementary Fig. 4). Increased mosaicity was observed by DFXM within the cycled Cu grains' volumes (Figs. 8,9) and by EBSD at the cycled film's cross-section (Fig. 7). However, DFXM indicated also that the microstructural refinement is more pronounced in the vicinity of HAGBs (Figs. 8a,b and 9a,b). The regions near HAGBs exhibit a higher mosaicity, which is irregular along the grains' circumference. This inhomogeneous grain refinement was interestingly not visible in the EBSD micrograph in Fig. 7, probably because EBSD probes only the sample's surface, while DFXM provides volume-averaged information. Additionally, the cyclic deformation resulted also in an accompanying accumulation of structural damage and also in the increase of residual elastic strain concentrations near HAGBs, as indicated by the FWHM and 2nd order X-ray elastic strain data, respectively, visible in Figs. 8c,d and 9c,d. Again, the increased elastic strain and structural defect densities in the vicinity of HAGBs occur locally and their occurrence does not always correlate completely. In contrast, the grains' interiors exhibit comparably lower 2nd order strain levels and defect concentrations as well as smaller mosaicity. In summary, DFXM indicates a different mode of thermomechanical damage in the vicinity of HAGBs compared to the grains' interiors. We interpret this as hardening of the regions near HAGBs [14,17,36,38].

Now the question arises if and how the hardened near HAGBs regions contribute to the observed voiding (cf. Figs. 8-9, supplementary Video 1). The brittle fracture of Cu GBs was reported already by Wimmer et al. [40] for micro-mechanically loaded Cu films tested at ~100 °C. The cracking was attributed to the deposition of S and Cl impurities at GBs, which stemmed from the electrodeposition process applied also in the present case. It can be therefore supposed that the early stages of the void formation in the present films (prepared using the same method as in Ref. [40]) can be also correlated with the occurrence of impurities. Next, our ex-situ EBSD data from a Cu thin film cross-section after 10⁴ thermal cycles show the presence of a near-interface void (solid arrows in Fig. 7a). This observation complement the data from Kleinbichler et al. [22], who reported a formation of voids exclusively at HAGBs and triple points.

Additionally, the present SEM, EBSD and X-ray nanotomography data (Figs. 1b, 6, 8e, 9e, 10 and Supplementary Video 1) document that the application of an even higher number of thermal cycles results in localized decohesion of HAGBs and in the growth of microscopic voids (Fig. 6). Up to now, the maximal number of applied thermal pulses of 1.2 × 10⁴, with a length of 200 µs, was reported by Moser et al. [21], who observed distinct surface roughening, intergranular voiding and deterioration of electric conductivity. The micrographs in Figs. 6 and 10 document that by increasing the number of applied thermal cycles to 5 × 10⁴, pronounced cross-sectional voids with a "width" of up to 2 µm are generated in Cu films.

A mechanism of void nucleation and their gradual growth at HAGBs of Cu films was reported by Schwaiger et al. [37]. During their cycling mechanical tests on Cu films, the alternating compressive and tensile stresses (Fig. 2b) initiated plastic deformation within Cu grains. In coarse-grained polycrystals, the dislocation motion is expected to be hindered by GBs, at which dislocations pile-up [16,24,38,37], which results in localized hardening [14,36]. Two opposite edge dislocations moving on adjacent slip planes may however annihilate [6,37], which may result in the formation vacancies especially at HAGBs and at surfaces, thereby acting as dislocation sinks. When the concentrations of vacancies becomes supersaturated, microscopic voids are formed at HAGBs, whose concentration and size increase with the number of

introduced vacancies [37]. The present data document (i) that the regions near HAGBs possess indeed a higher amount of structural defects compared to the grain interior and (ii) that these regions are expected to be hardened due to the observed microstructural refinement, which induces also higher elastic strain levels (Figs. 8,9). Therefore, it can be expected that the observed microstructural refinement of regions near HAGB promotes the decohesion of HAGBs (Fig. 6 and Supplementary Video 1) and results in their fracture.

Methodologically, the DFXM approach [26] allows for the first time to analyze microstructural damage in thermally cycled Cu thin films after more than 10³ thermal cycles and to assess microstructural changes not only on the film's surface but instead to provide film thickness-averaged information. In all of the previous studies, the application of more than 10³ thermomechanical cycles resulted in surface roughening of the films, which prevented EBSD analysis. Moreover, compared to previous studies based mostly on SEM, EBSD and transmission electron microscopy, the DFXM approach allowed us to quantify also (projections of) intragranular elastic strains of the 2nd order across grains, which were found to be correlated with the FWHM data. This unique correlation provides an opportunity to understand the relationship between microscale intragranular elastic strain distributions and their microstructural origins in metals after large accumulated deformations, such as severe fatigue experiments.

## 5. Conclusions

The main aim of this study was to analyze the impact and origins of the thermomechanical fatigue within 20 µm thick Cu thin films after 5 × 10⁴ thermal cycles in the temperature range of 100–400 °C using a heating rate of 10⁶ K/s. The application of ex-situ DFXM, ex-situ EBSD and in-situ SEM revealed microstructural refinement of individual Cu grains, hardening in regions near HAGBs and possible origins of HAGB decohesion and voiding.

The main results can be summarized as follows:

- Thermal cycling of the Cu film resulted in severe shear deformation of Cu grains and accumulation of plastic strain, which is accompanied by substantial film surface roughening, the formation of LAGBs within the grains and intergranular "micro voids" at HAGBs (Figs. 1b,6,7).

- During the comparatively fast thermal cycling with 5 × 10⁴ thermal cycles, through-film-thickness voids with a width of up to 2 µm developed gradually, preferably at HAGBs (Figs. 6,10, Supplementary Video 1).

- Ex-situ DFXM indicated an increased concentrations of tensile and compressive residual elastic strains of the 2nd order and structural defects near HAGBs, whose occurrence correlates well with elevated mosaicity (Figs. 8,9). However, the grains' interiors exhibit moderate strain levels, defect concentrations and mosaicity.

- Based on the published literature data and models, the observed decohesion of HAGBs was interpreted by the hardening of near-HAGB regions, as a result of vacancy condensation in front of the HAGBs, after dislocations have moved across cycled grains and partly annihilated.

Methodologically, the work documents the advantages of DFXM, which allows to analyze intragranular mosaicity, strains and defect densities in bulk Cu grains. As an outlook, it can be expected that the already finished EBS update at ESRF in Grenoble will allow for in-situ and operando experiments, enabling us to monitor intragranular dynamics of Cu grains during thermal cycling.

## Declaration of Competing Interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence
the work reported in this paper.

## Acknowledgment

JK is grateful to Prof. Daniel Kiener for valuable discussions on Cu twins. The authors gratefully acknowledge the financial support under the scope of the COMET program within the K2 Center "Integrated Computational Material, Process and Product Engineering (IC-MPPE)" (Project No. 859480). This program is supported by the Austrian Federal Ministries for Transport, Innovation and Technology (BMVIT) and for Digital and Economic Affairs (BMDW), represented by the Austrian research funding association (FFG), and the federal states of Styria, Upper Austria and Tyrol. JK would like to thank the Research Infrastructure project SOLID21 (project No. CZ.02.1.01/0.0/0.0/16_019/0000760) and CzechNanoLab (LM2023051) supported by MEYS CR. This work was funded by the Austrian Research Promotion Agency (FFG, Project No. 884573). The authors acknowledge the ESRF for the provision of beamtime on the beamline ID06-HXM. SF gratefully acknowledges financial support from the Deutsche Forschungsgemeinschaft [DFG, German Research Foundation, SFB 986 (project Z2)]. We thank Berit Zeller-Plumhoff for assistance with the 3D visualization of the nanotomography data.

## Supplementary materials

Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.actamat.2023.118961.

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