2.1.

Laser-Induced Surface Modification

The leading CdTe solar cells are based on a superstrate design (glass/TCO/CdS/CdTe/back contact), where the CdTe absorber is the final layer deposited before back contact metallization with the light impinging through the glass substrate. Ohmic contacts are vital in solar cells as they allow for the maximum collection of photogenerated carriers. The formation of an Ohmic contact in the final metallization step has long been a difficult task due to CdTe’s high electron affinity and its propensity to form surface dipoles that pin the Fermi level.²⁵ As a result of these dipoles, even very large work function metals do not necessarily make an Ohmic junction.²⁶ A non-Ohmic contact in a solar cell behaves essentially as a reverse polarity Shottky diode and results in a second depletion width in the cell leading to poor fill factor and losses in open-circuit voltage ($V_{\mathrm{oc}}$) and short-circuit current ($J_{\mathrm{sc}}$),²⁷ all of which lower the device’s efficiency. Several engineering techniques have been devised to create quasi-Ohmic contacts including the addition of highly doped p-type buffer layers, preferential chemical etches, or combinations of both. A heavily p-doped surface region between lightly p-type CdTe and metal will create an Ohmic junction by allowing hole flow through a narrowed tunneling region. Recently, PLA has been used for creating an Ohmic back contact.^28,29 A pulsed laser-based approach has several advantages including not requiring additional vacuum processing, being highly tunable through pulse number and fluence, and it does not create a Cd-contaminated liquid waste stream as is the case with chemical etches.

An Ohmic back contact is created with PLA by exploiting the difference in vapor pressure between Cd and Te, with Cd being the more volatile species.^30,31 When a pulsed UV laser is applied, the surface rapidly heats to very high temperatures causing slightly more Cd loss than Te. This can leave either a Te-rich surface, which effectively dopes the surface highly p-type due to the increased presence of Cd vacancy acceptors, or leaves a thin Te film, or some combination of both.³² Elemental Te is a degenerate p-type semiconductor and thin layers on the surface of CdTe have been shown to lead to Ohmic behavior,³³ although this may be due more to the formation of metal tellurides than to the properties of Te itself, as has been discussed.³⁴ Pulsed laser-induced stoichiometry changes were first noted, albeit casually, by Montgomery et al. as they investigated the use of 400 ns pulses of 590 nm light to prepare clean crystalline CdTe surfaces.³⁵ Brewer et al. more fully explored the phenomenon using a KrF excimer laser with 15 ns pulses and showed that above a threshold fluence, a Te-rich surface was created before the onset of melting.³¹ Furthermore, it was discovered that this behavior was completely reversible by the application of sub-threshold light, which suggested to the authors that the mechanisms responsible were a temperature-related competition between Cd and ${\mathrm{Te}}_2$ desorption. A Te surface enrichment with ruby or KrF laser pulses has been experimentally confirmed by several other groups.^36–38 Golovan et al. calculated the surface temperature reached during PLA by accounting for the heat losses due to evaporation of Cd and Te.³⁹ This was later adapted by us for use in a three-dimensional finite-element model of laser interactions with CdTe.¹⁸

Pulsed laser-induced surface modification was first applied to the more PV-relevant polycrystalline CdTe by Nelson et al. by annealing 200 to 300 nm thick CdTe films using 8 ns pulses of the second, third and fourth harmonic of a Nd:YAG laser.⁴⁰ They performed x-ray photoemission microscopy to show that a Te-rich surface was produced. These films were highly idealized as they were grown in ultra-high vacuum and had a much larger grain size to film thickness ratio ($>1\mu \mathrm{m}$: 200 to 300 nm) than is typical of PV device-quality CdTe. Simonds et al. showed that PLA using a KrF excimer laser of thin film CdTe was capable of making an Ohmic back contact in CdTe PV devices by applying laser pulses before back contact metallization.²⁸ Transfer length measurements showed that a single, 25-ns pulse was capable of lowering the specific contact resistivity by a factor of over 20. Current-voltage measurements of completed devices revealed that after 100 pulses, all signs of a non-Ohmic back contact had disappeared. Further investigations of the surface chemistry induced by PLA suggested that the Te-layer was not uniform but concentrated at the exposed surfaces at grain boundaries,³² which was also noted previously.⁴⁰ Although it is clear that a PLA is capable of much more rapidly producing an Ohmic back contact, further work is necessary to show how the device performance compares with other established techniques of creating Ohmic contacts. This is an ongoing topic of our research.

During standard deposition processes of thin film CdTe for PV, it is empirically well-known that CdTe congruently sublimes, meaning stoichiometric films are readily attained. However, the above literature shows that with PLA, this is not the case. In fact, even CW illumination has been seen to cause incongruent sublimation of Cd and Te.^15,41–45 This was demonstrated by Soares et al. using micro-Raman spectroscopy, where a distinct spectral line from elemental Te was seen to grow as a result of increasing intensity of the Raman excitation laser source.^44,45 An example of this phenomenon is shown in Fig. 5 from our own micro-Raman measurements where an as-deposited polycrystalline CdTe sample is continuously exposed to the excitation light with spectra taken every 5 min. One sees that a spectral line (${\mathrm{A}}_1$) due to an optical mode of elemental Te⁴⁶ near $120{\mathrm{cm}}^{-1}$ grows as a result of CW exposure to the excitation source. Soares et al. were able to calculate the surface temperature by comparing the intensity ratios of the Stokes and anti-Stokes Raman signals and found that in their measurement, the surface was $\sim 200°\mathrm{C}$ above room temperature, which is well below the melting threshold of 1092°C for CdTe.⁴⁵ A careful measurement of the crater depth versus temperature created by CW illumination is compared to that expected by purely thermal sublimation of Cd and ${\mathrm{Te}}_2$ by Uzan et al., who revealed that the activation energy of the laser induced process was 1.44 eV smaller than the assumed thermal process.¹⁵ The fact that this value is nearly identical to the bandgap of CdTe (especially at elevated temperature) suggests that a recombination-induced athermal mechanism is responsible for Cd vacancy, ${\mathrm{V}}_{\mathrm{Cd}}$, and creation. Furthermore, Tsyutsura and Shkumbatyuk showed that it was possible to use CW ${\mathrm{CO}}_2$ laser light to create a shallow p-n junction in indium-doped n-type crystalline CdTe.⁴⁷ They claimed that the junction formation was due to a laser-induced redistribution of n-type ${\mathrm{V}}_{\mathrm{Cd}}\text{-}{\mathrm{In}}_{\mathrm{Cd}}$ complexes into In interstitials and ${\mathrm{V}}_{\mathrm{Cd}}$ acceptors. In light of this evidence, it seems most likely that additional ${\mathrm{V}}_{\mathrm{Cd}}$ were also being produced during CW LA thereby being at least partially responsible for the creation of the p-type region. These hypotheses have not been completely explored but have shown that LA allows for processing conditions that are not possible by traditional thermal means.

Fig. 5

Micro-Raman measurements at 5-min intervals while being continuously exposed to the 532-nm CW excitation source. The peak near $120{\mathrm{cm}}^{-1}$ is due to elemental Te (${\mathrm{A}}_1$). The broad peak at $145{\mathrm{cm}}^{-1}$ is an overlap of a Te peak (${\mathrm{E}}_1$) and a transverse optical (TO) mode of CdTe. The low intensity shoulder near $165{\mathrm{cm}}^{-1}$ is due to CdTe.

2.2.

Laser Processing for Bulk Effects in CdTe

Recombination of photoexcited carriers by deep-defects results in very short minority carrier lifetimes and low open-circuit voltages and is, therefore, a major culprit in limiting the performance of CdTe devices.⁴⁸ Additionally, the deposition of high quality polycrystalline CdTe is currently one of the slowest steps in completing a solar cell module and requires process parallelization in order to maintain the total average cycle time of the other, much faster, module production processes.⁴⁹ Therefore, a method of rapidly depositing CdTe, even if initially lower in quality, followed by a rapid LA process to increase grain size and anneal defects could potentially save on capital costs and increase module throughput. In this section, we will review evidence from literature which shows that LA is capable of achieving both of these material effects.

First, as a proof-of-principle that LA can be used to rapidly improve initially low-quality material, Fig. 6 shows x-ray diffraction (XRD) data before and after CW 1064 nm LA for 100 s of sputtered CdTe made without substrate heating during deposition. Although the laser photon energy (1.17 eV) is below the nominal bandgap of CdTe, appreciable heat is generated due to direct absorption by defect states.¹⁷ Four distinct peaks are seen: (100), (101), and (002) from a wurtzite phase⁵⁰ and the {111} zincblende phase with the latter two overlapping. The presence of a local wurtzite crystal structure is synonymous with stacking faults which result in poor electrical quality.⁵¹ Figure 6 clearly shows the ability of LA to reduce the presence of the wurtzite phase and the extended defects associated with it: the wurtzite peaks decrease in intensity and the zincblende peaks become narrower upon LA.

Fig. 6

XRD data of as-deposited, low quality CdTe that was laser annealed with CW 1064 nm light with 100 s dwell time.

The LA of polycrystalline CdTe was first studied by Dawar et al., who used a ns-pulsed 1064 nm Nd:YAG laser on 700 nm CdTe.⁵² By increasing the pulse energy and number, they showed that grain growth from 45 to 103 nm could be achieved, as well as a more than doubling of the mobility. They also measured, via Hall measurements, an increase in the carrier activation energy that they attributed to defect creation, which increased with pulse fluence and number. This is perhaps a result of stress introduced during PLA and is possibly why subsequent studies have relied on CW laser anneals. Indeed, our own initial findings with PLA of defective CdTe showed grain growth, but increased levels of strain are not shown. Grain growth has also been seen with CW LA with an 808-nm light source rapidly scanned across the film surface.⁵³

Laser recrystallization of large-grained epitaxially grown films has also been shown to quickly increase grain size. As and Palmetshofer⁵⁴ showed that the maximum grain size increased after only 2 s exposure of CdTe:In films to 647-nm CW light. Defect concentrations were compared between laser annealed and isochronal thermal annealed samples with photoluminescence (PL) and deep-level transient spectroscopy. These measurements revealed that the laser and thermal annealing mechanisms are fundamentally different in that deep defects, not seen after LA, bet appear after thermal annealing at comparable temperatures. Other LA studies of epitaxial CdTe showed that LA was capable of relieving microstresses in the films as well as decreasing surface roughness.⁵⁵ LA has also been investigated to repair radiation-induced damage,^56–59 where in some cases, it has been found to be more effective than furnace annealing.^57,59 For instance, Schuller et al. found that a 2-s CW LA gave similar lattice repair to radiation damaged CdTe as a 15 min isothermal furnace annealing.⁵⁹

Recently, our group has shown that sub-bandgap CW light (1064 nm, Nd:YAG) can improve the optical and crystalline quality of initially low-grade polycrystalline CdTe.¹⁷ The absorption in these films is due to defect states, which were shown to be reduced upon LA. This was monitored in situ by measuring the partially transmitted laser light during the annealing. This work has not only shown the ability of LA to rapidly improve material quality, but also revealed a new technique by which real-time process control can be implemented to improve throughput.

Several novel methods of using lasers to rapidly crystallize precursor films into thin film CdTe have been studied. Baufey et al. deposited alternating layers of Cd and Te and employed pulsed and CW laser light to synthesize and crystallize thin film CdTe.⁶⁰ They found that the LA parameters can be varied to successfully obtain PV-relevant zincblende phase CdTe with 500-nm sized grains. The CW annealing gave larger grain sizes that were single zincblende phase, whereas the PLA resulted in a mixed phase that included wurtzite crystal structure. Electrophoresis using polar organic solvents followed by LA has also been explored as a low-cost deposition route and has achieved deposition rates of $20\mu \mathrm{m}/\min$.⁶¹ A rapidly scanned (0.2 to $0.6\mathrm{mm}/\mathrm{s}$) argon laser line (514 nm) produced the highest quality films when compared to furnace or PLA with a 1064 nm light. They also found that the presence of ${\mathrm{CdCl}}_2$ was vital for enhancing recrystallization. Additionally, pulsed UV lasers have been used to photodissociate organometallic precursors to create epitaxial CdTe thin films.^62,63 However, the very slow deposition rates on the order of $1\mu \mathrm{m}/\mathrm{h}$ will ultimately limit the usefulness of this technique for PV applications.

A more promising route for using LA to rapidly produce thin film CdTe from a low-cost precursor is from the sintering of CdTe nanoparticles that can readily and cheaply be chemically synthesized. Rickey et al. employed a novel technique of pulsed-laser induced compression followed by PLA.⁶⁴ Several hundred MPa of compression were first applied to nanoparticle films by an unspecified “highly absorbing material” in intimate contact with the nanoparticle mesh that was rapidly expanded by the absorption of a 5 ns, 1064 nm laser pulse. The compressing material was then removed and two 7.05 mJ, 25 ns pulses (wavelength not specified) were applied to the compressed film. It was found that these two pulses were sufficient to sinter the particles and achieve film resistances similar to those found in conventional polycrystalline CdTe. Unfortunately, no optoelectronic characterization was undertaken and the films were far from optimal. However, this or related techniques seem to warrant further investigation for low-cost PV applications.

3.1.

CW Laser Annealing

Figures 7a(i) and 7a(iii) depict the two precursor types that have been used for CW LA. The annealing objectives in both cases are to fully homogenize the elements and promote $\mathrm{Cu}(\mathrm{In},\mathrm{Ga}){\mathrm{Se}}_2$ grain growth. The first precursor, shown in Figure 7a(i), consists of the constituent elements deposited by electrodeposition in a single step at room temperature to form a nanocrystalline film. While the nanocrystallinity leads to high excess grain boundary energy, there is only a minimal chemical driving force to change the precursor since the precursor consists of binary and ternary selenides. By contrast, the metallic elements, Cu, In, and Se, can also be deposited in separate steps to form a stack, as shown in Fig. 7a(iii). This second precursor type has a much greater requirement for atomic diffusion and chemical reaction than the co-deposited precursor, but is chemically much more reactive since the elements have an oxidation state of zero. In order to stimulate the transformation processes of atomic diffusion, grain growth and chemical reaction, the LA must supply sufficient energy. Below, the approaches are described which have, to date, been used to achieve this aim.

3.1.2.

Annealing of co-deposited Cu-In-Se

Subsequent approaches to the CW LA of CISe have focused on co-deposited Cu-In-Se films [see Fig. 7a(i)], where stoichiometric quantities of each of the constituent elements are deposited in a single step. This process distributes Se evenly through the film depth and the majority is bound to Cu and/or In atoms as opposed to existing in its free state. Jost et al.⁷⁵ initiated this processing route by using a Nd:YAG ($\lambda =1064\mathrm{nm}$) to anneal co-electrodeposited CISe films. Comparing the results of samples annealed at incident laser fluxes from 11.6 to $36.8\mathrm{kW}{\mathrm{cm}}^{-2}$ to those that had been annealed in a furnace for 5 min at 550°C, the LA absorbers demonstrated higher crystallinity (as determined from the narrowing of CISe Bragg-peaks in XRD diffractograms). This result was achieved by subjecting the film to rapid thermal cycling achieved by scanning the $110\mu \mathrm{m}$ beam at $100\mathrm{mm}{\mathrm{s}}^{-1}$ 20 times. This gave a total annealing time of 22 ms, which is four orders of magnitude faster than “rapid” annealing processes. However, as with the rapid annealing of SELs in the previous section, the CISe film was seen in SEM micrographs to dewet the substrate.

Consequently, our work on CW annealing of CISe focuses on solid state processing. This was initially achieved by a significant drop in laser flux to $0.05\mathrm{kW}{\mathrm{cm}}^{-2}$ to avoid melting, with typical anneal times between 15 and 60 s.⁶⁷ The SEM of such CISe films’ surfaces showed no change in their surface appearance with LA, consistent with a non-melt process. Analyzing the bulk of the material using XRD showed that LA increased the crystalline order determined by a narrowing of the characteristic chalcopyrite peaks in the x-ray diffractogram. The effect of metal composition on the crystallinity was tested by annealing “Cu-rich” ($\mathrm{Cu}/\mathrm{In}>1$) and “Cu-poor” ($\mathrm{Cu}/\mathrm{In}<1$) precursors. The CISe crystal structure is tolerant to a wide range of stoichiometries⁷⁶ and the adjustment of this parameter remains an interesting research topic.⁷⁷ While record CISe devices have a Cu-poor final composition,⁷⁸ they often include a Cu-rich growth step that is observed to enhance optoelectronic and structural properties.⁷⁹ Our group⁶⁷ also measured a narrower 112 XRD peak for laser annealed absorbers from Cu-rich precursors when compared to their Cu-poor counterparts. Mooney and Hermann⁸⁰ revealed that the extinction coefficient, $k$, which describes the attenuation of electromagnetic radiation by a material, is dependent on stoichiometry and ranged from 0.15 for Cu-poor to 0.3 for Cu-rich electrodeposited precursors. A higher $k$ leads to a higher power density being dissipated in the absorber, as demonstrated in the calculations of Ref. 67, which used a transfer matrix approach to analyze power absorption versus depth. The increased energy absorbed by Cu-rich films results in higher peak film temperatures during LA, and a faster rate of atomic diffusion, grain growth and defect removal than the Cu-poor samples. So, both higher optical absorption and greater atomic mobility are related to Cu richness and probably combined to result in the observed differences versus precursor film Cu composition.

Initial work by our group established a 15 s minimum annealing time to yield structural improvements in electrodeposited CISe films.^67,81 Although 15 s is faster than a 5-min RTP process⁷⁵ or up to 60-min furnace process,⁸² it neither demonstrates the full potential of this rapid heating technique nor offers a significant energy savings over current industrial furnace processing. Furthermore, shorter annealing times, particularly when achieved with a technique compatible with roll-to-roll processing, can dramatically increase industrial throughput. Consequently, further work has focused on an “intermediate” time/flux regime.^83–86 Fluxes between 150 and $945\mathrm{W}{\mathrm{cm}}^{-2}$ and a dwell time up to 1 s gave a noticeable structural improvement. This is measurable as an observed increase in grain size between precursor and annealed samples in SEM [see Figs. 7a(i) and 7a(ii)] and the CISe peaks in the XRD diffractogram become narrower with a longer annealing time [illustrated in Fig. 8(a)]. Calculating the crystal coherence length from these peaks using the Scherrer equation shows apparent average crystallite domain sizes of 43 to 140 nm for 0.25 to 1 s dwell time [Fig. 8(b)]. The 140-nm crystallite domain size achieved with 1-s dwell time, and $150\mathrm{W}{\mathrm{cm}}^{-2}$ is larger than that achieved for $50\mathrm{W}{\mathrm{cm}}^{-2}$ for 15 to 60 s.⁶⁷ Hence, crystallite domain size growth is a function of both flux and time, and as will be seen below, also depends on the selenium vapor pressure above the sample.

Fig. 8

(a) XRD diffractograms ($\theta -2\theta$) of ${\mathrm{CuInSe}}_2$ absorbers focusing on 112 peak, after LA Cu-In-Se precursors with Nd:YAG laser at $150\mathrm{W}{\mathrm{cm}}^{-2}$, with scan speeds (corresponding to annealing times) of (i) $2\mathrm{mm}{\mathrm{s}}^{-1}$ (1 s), (ii) $4\mathrm{mm}{\mathrm{s}}^{-1}$ (0.5 s), (iii) $6\mathrm{mm}{\mathrm{s}}^{-1}$ (0.3 s), (iv) $8\mathrm{mm}{\mathrm{s}}^{-1}$ (0.25 s), (beam area: $2\times 2{\mathrm{mm}}^2$), compared to (v) precursor. (b) Crystallite coherence lengths of determined from FWHM of 112 peak of the diffractograms of (a).

Importantly, this work was the first to test the optoelectronic properties of an absorber formed by LA, which indicated an increase in photoelectrochemical photocurrent and a higher level of radiative recombination as compared to the precursor. However, the initial device efficiency from an absorber formed in this process was negligible. An XRD study,⁸⁴ evidenced that these absorbers may contain an unreacted ${\mathrm{In}}_2{\mathrm{Se}}_3$ secondary phase, and/or structural defects, quenched by the rapid cooling following laser irradiation. The cause for the low efficiency, however, was believed to stem from the loss of volatile Se from the film, even within the 1-s annealing time.⁸³ Se may be lost directly from the Se lattice resulting in ${\mathrm{V}}_{\mathrm{Se}}$ defects, which have been correlated to poor device performance,⁸⁷ or resulting from the decomposition of the CISe semiconductor into its binary selenides.

In order to prevent Se loss, subsequent LA experiments were carried out under a background partial pressure of Se.^85,86,88 A higher Se activity led to an increased $\mathrm{Se}/(\mathrm{Cu}+\mathrm{In})$ ratio in the absorber, as well as larger grains and a higher PL yield as compared to lower Se activity.⁸⁶ Figure 9(a) illustrates the difference in the PL yield of absorbers annealed under high and low partial pressures of Se. Implementing this new approach, the first and only known working solar cell device to date with a LA absorber was fabricated with a power conversion efficiency of 1.6%.⁸⁸ Characteristic patterns on the surface of this absorber [Fig. 8(b)] appear to be linked to the rastering of the $2\times 2{\mathrm{mm}}^2$ laser beam over the surface. The Gaussian spatial profile of the power distribution was shown to affect the annealing temperature profile of the CISe film.⁸⁵ This uneven energy distribution stimulates different levels of grain growth and atomic diffusion perpendicular to the beam scanning direction, and consequently, variations in lateral optoelectronic properties are expected. As is well known, fluctuations in properties for solar cells result in voltage and efficiency loss; essentially, the cell performance is linked to the lowest quality material in the device. It is believed that with a homogeneous beam flux, and minimal thermal variance over the sample area, a consistent photon conversion would be achieved over the full device area, and higher power conversion efficiencies would be achieved.

Fig. 9

(a) PL Spectra of ${\mathrm{CuInSe}}_2$ co-deposited precursor and after annealing at $150\mathrm{W}{\mathrm{cm}}^{-2}$, 1 s dwell time per unit area, with (ii) a low PSe ($\approx 9\times {10}^{-7}\mathrm{Pa}$), and (iii) a higher PSe ($\approx 2\times {10}^{-4}\mathrm{Pa}$). These areas correlate, respectively, to the pale and dark lines observed on the absorber surface as indicated in (b).⁸⁸

LA of co-deposited precursors shows that it is possible to stimulate sufficient atomic diffusion, chemical reaction, grain growth, and enhancement of optoelectronic properties with $\le 1\mathrm{s}$ dwell time while remaining in the solid state regime. This process has enabled the first device to be fabricated where transformation from CISe precursor to absorber was achieved using only a laser. Although the efficiency of this device is low, key points for improvement are identified, which include homogenizing the laser beam flux and increasing the background Se activity during annealing. This provides opportunity for future work to increase device efficiency and consequently allow this processing route to compete with current methods.

3.2.

Pulsed Laser Annealing

An alternative LA process tested on CI(G)Se is PLA. This differs substantially in both the material and annealing objectives as compared with CW LA. CW annealing gives a constant power output, whereas PLA is characterized by high peak power but a much lower average power. The energy supplied during the very short pulses will rapidly heat the film. However, the pulse times remain below those required for thermal equilibrium, leading to large temperature gradients. PLA initially found favor in removing damage caused by ion-implantation in the near-surface regions of Si wafers.⁸⁹ PLA of CI(G)Se aims to modify defects and thus enhances final device efficiency. All reported studies of PLA in CI(G)Se use a 248-nm excimer laser (pulse width, $w\approx 25\mathrm{ns}$), which has a penetration depth of 50 nm in CISe and a heat affected region [Eq. (1)] of about 300 nm with $D$ estimated from thermal conductivity, density, and specific heat values.^11,90 Therefore, only the near surface region of the CI(G)Se film is affected by the laser energy, with the film bulk remaining unaltered (see Fig. 3). As a result of $l_t$ being much less than the film thickness, the bulk of the film must already show suitable characteristics for a device prior to annealing, therefore, these studies use only “high” quality CISe/CIGSe absorbers, formed using either method (1) single step co-evaporation, or a precursor formed using method (2) that has been thermally treated prior to LA.

The unsuitability of PLA for absorber formation is exemplified in the work of Bhatia et al., who carried out PLA on similar electrodeposited CISe precursors to those which were successfully CW LA using a 1064-nm Nd:YAG laser. While both the Nd:YAG laser in pulsed mode (pulse time 75 to 150 ns) and a 248 nm KrF laser ($\approx 25\mathrm{ns}$) induced some crystal growth at high flux, they also caused the CISe to melt and dewet.^81,91,92 PLA in the sub-melting regime ($\le 70\mathrm{mJ}{\mathrm{cm}}^{-2}$) had no measurable effect on the structural properties of the CISe film. As this fluence exceeded that which was later shown to produce both structural and optoelectronic improvements in high quality CISe absorbers, it is concluded that this result is due to the poor thermal transport in the nanocrystalline and multiphase, as-deposited CISe precursors.⁸³ Only by melting the film is a sufficient atomic mobility and energy available for CISe formation throughout its full thickness. As the dewetting of the liquid phase CISe occurs spontaneously on all tested substrates,⁶⁷ it was concluded that there is no practical processing window that allows PLA to be used for CISe absorber formation.

The first published studies of sub-melt PLA of CISe were carried out by the group of Anderson,^93–95 which used co-evaporated $\mathrm{Cu}(\mathrm{In},\mathrm{Ga}){\mathrm{Se}}_2$ absorbers as pictured in Fig. 7b(i) with a CdS buffer layer deposited on the surface. Having identified that interfacial recombination at the CIGSe/CdS interface is a limiting factor to optimal device performance the authors used PLA to modify near-surface defects in the CIGSe absorber. Initially, they tested the effect of five pulses with energy densities from 20 to $60\mathrm{mJ}{\mathrm{cm}}^{-2}$. XRD showed a slight increase in crystallographic order compared to the un-irradiated film, indicating that the laser energy is sufficient to stimulate atomic diffusion. All conditions were shown by dual beam optical modulation to increase the charge carrier lifetime, with the lowest fluence giving the most significant improvement with a near tripling of this value. Furthermore, Hall measurements showed an increase in carrier mobility and decrease in film resistance. Both of these results are consistent with PLA removing near-surface electrically active defect states. As these states act as carrier traps and recombination centers, it was expected that PLA should lead to devices with higher efficiencies. However, devices formed from samples irradiated at $50\mathrm{mJ}{\mathrm{cm}}^{-2}$ showed equivalent efficiencies to the un-irradiated control. This was explained using external quantum efficiency measurements, which showed a decrease in carrier collection at wavelengths $<630\mathrm{nm}$ after laser irradiation. This indicated that the poor device performance could be due to the laser pulse also causing damage in the CIGSe/CdS interface region. Increasing the pulse number to 10 or 20 pulses⁹³ also showed similar improvements in Hall mobility. However, device measurements revealed that only the lowest laser fluence ($30\mathrm{mJ}{\mathrm{cm}}^{-2}$) led to an efficiency increase over the 9.1% efficient reference device. Fluences of $40\mathrm{mJ}{\mathrm{cm}}^{-2}$ or higher degraded the cell performance, consistent with increased laser damage. Measurements of the low-fluence annealed devices yielded 12.1% and 11.1% efficiencies for 5 and 10 pulses, respectively. Although not a recorded result, it does show a remarkable 33% relative improvement with a process that takes less than 1 s to perform.

Ahmed et al. carried out a detailed analysis of the defect energy levels in CISe films formed using flash evaporation.⁹⁶ The slight increase in bandgap from 1.00 to 1.02 after five pulses ($w\approx 20\mathrm{ns}$ time) of laser irradiation at 170 mJ was claimed to be related to a modification of defect levels at the band edge. High-resolution photoacoustic spectroscopy revealed the ionization and proposed electrical activities of five deep-level defects, which provide non-radiative recombination pathways within the film. However, upon close scrutiny, it is clear that these purported defect absorption transitions correspond to the narrow peaks in the illuminating lamp’s spectrum. It is well known that it is difficult to normalize such spectra, and it is common to see artifact peaks arising from imperfect normalization. Also, defect-to-band transitions, which should be more probable than discrete defect-to-defect transitions, will produce an onset in absorption and not a resonant peak. We conclude that it is possible that some changes in sub-bandgap absorption were responsible for the observed changes in photoacoustic absorption; however, the data are rife with experimental artifacts so the conclusions should be carefully examined. The position and intensity of near-edge absorption features were claimed to become less pronounced with LA but with an additional consequence of forming a shallow defect near the band edge.

Bhatia et al. measured the PL yield after LA of co-electrodeposited CISe samples with five pulses of fluences from 20 to $40\mathrm{mJ}{\mathrm{cm}}^{-2}$.^91,97 Current-voltage curves measured in the dark showed an improvement in the diode ideality factor at fluences of 20 and $30\mathrm{mJ}{\mathrm{cm}}^{-2}$.⁹⁷ The increased PL yield with LA was attributed to an increase in radiative transitions due to a reduction in deep-defect concentration. However, an observed PL broadening was indicative of a more highly compensated semiconductor, which can be caused by an increased concentration of shallow defects. Deep-level transient spectroscopy and temperature-dependent admittance spectroscopy indicated that this shallow trap state was around 60 meV and detected the presence of a majority hole trap at 200 meV. The formation of shallow defects may be due to laser processing in an inert atmosphere. A low Se activity during annealing has been shown to lead to formation of ${\mathrm{V}}_{\mathrm{Se}}$ and the shallow defect ${\mathrm{In}}_{\mathrm{Cu}}$, both of which are detrimental to device performance.⁸⁷ This assumption is further supported by Ref. 98 where the Schottky barrier height for the CISe film annealed at the highest fluence indicated the presence of oxide phases and/or a high density of interface states, possibly from Se vacancies.

In this section, it has been shown that pulsed, non-melted LA has the potential to improve the structural and optoelectronic characteristics of chalcogenide absorber layers. By reducing the deep-level defects, particularly those localized at the film surface, a material with an improved carrier lifetime is achieved. This has the potential to increase the short circuit current, $J_{\mathrm{Sc}}$, of the final devices fabricated from these films, and therefore, increase device efficiency. However, a counter effect of the technique appears to be the creation of shallow defects and surface damage with higher fluences. Consequently, there is only a narrow processing window for PLA of CI(G)Se, in which the potential of this processing step is realized.