FASES DEL PROYECTO DE INVESTIGACIÓN DE TIPO DESCRIPTIVA.

With the continued investigation of aluminum anodes for lithium (primary) batteries interest eventually grew for using this material to replace carbon in lithium ion (secondary) batteries. Based on the stoichiometry and the low atomic weight of Al the LiAl, Al2Li3 and Al4Li9 alloys offer theoretical mass capacities of 993, 1490 and 2234 mAh/g respectively. Therefore the maximum theoretical lithium uptake for an aluminum electrode is 2.25 Li atoms per Al atom. This value is below the 4.25 or 3.75 Li atoms for each Sn or Si atom possible in the highest order Li17Sn4 and Li15Si4 alloys at room temperature [4]. However as described in Ch. 2.2 both Sn and Si intermetallic alloy formation suffer from severe volumetric expansions of 676% and 323% respectively for the highest order alloys which causes rapid irreversible capacity loss as anode materials [3]. Studies of Al anodes in lithium ion batteries over the past 15 years have revealed that only the lowest order (LiAl) alloy phase is formed during lithiation-delithiation in aprotic polar carbonate solvents at room temperature [14-24]. This is consistent with the absence of higher order lithium-aluminum alloys at room temperature in primary lithium batteries discussed in Ch. 2.3.2. The structural determination of the intermetallic phase in Al anodes is typically made after galvanic cycling through X-ray diffraction (XRD).

While this lowest order intermetallic alloy only allows for an uptake of one Li atom per Al atom it still offers a good theoretical mass capacity of 993 mAh/g compared to the value of 372 mAh/g for graphite. Furthermore the volumetric expansion for LiAl alloy formation is only 97% [3], which is considerably less than for the higher order Si and Sn based alloys described above. The formation and dissolution of LiAl is a single alloy transition unlike the stepwise pathway for formation of higher order Si intermetallic alloys. Therefore the charge-discharge curves for Al anodes are characterized by wide and flat charge/discharge plateaus. This is an important requirement for steady power output of anodic materials in lithium ion batteries.

Recently the mechanism of lithiation in Al anodes has been investigated both in crystalline bulk materials and semi-crystalline nanostructures. For bulk Al materials Liu et al. have studied the mechanism of lithiation in thin Al foil, shown in Fig. 2-3-5 [25]. This figure is only a schematic representation of the processes and certain details such as solid-electrolyte interphase (SEI) formation and volume changes of the Al host material have been omitted. First the lithium ions arrive at the electrode surface (panel a). Lithiation of Al then proceeds through an initial nucleation of an α-LiAl solid solution in

the electrode surface (panels b,c). This process is accompanied by a lattice contraction in the Al host through rearrangement of domains. As the Li concentration in the Al matrix increases towards supersaturation a crystalline phase of β-LiAl begins to nucleate and

grow within the solid solution of α-LiAl (panel d). This process is accompanied by a

lattice expansion in the Al host. Therefore a two-phase equilibrium between α and β

phases of LiAl is established, leading to the lithiation potential plateau. As lithiation continues additional Li will diffuse into the bulk of the anode as a progressively deeper front of α-LiAl, while additional α-LiAl crystallizes into β-LiAl closer to the anode

surface (panel e inset). This electrochemical lithiation of Al through a solid-solution mediated pathway leading to crystallization is fundamentally different from the amorphization pathway described previously for other metal-alloying anode materials such as Si and SnO2 [4]. Physically the creation and movement of these LiAl phases is based on the mobility of dislocations within the Al host. Huang et al. revealed that initial nucleation of intermetallic phases in metal-alloying anodes occurs at surface sites with high density of mobile dislocations [26]. The accompanying lattice contractions and expansions of lithiation proceed through movement of these dislocations, creating a dislocation-induced stress (DIS) that is localized to the surface of the material during initial lithiation [27]. As lithiation continues (time increases) and the diffusion front moves deeper these dislocations will progressively move towards the interior of the anode bulk. As a result the DIS will decrease at the surface and increase in the bulk, eventually reaching a steady state.

Figure 2-3-5: Schematic of lithiation mechanism in a crystalline thin Al foil anode, from ref. 25.

A common choice for nanostructured anode materials is thin semi-crystalline films prepared on hard substrates through electron-beam deposition or thermal evaporation. Recently Leenheer et al. utilized in-situ transmission electron microscopy (TEM) to investigate the mechanism of lithiation in a 50 nm Al thin film prepared through electron- beam deposition on silicon nitride substrate [28]. Initially there were isolated nucleation events at the anode surface involving formation of α-LiAl solid solution, followed by

crystallized growth into the β-LiAl phase. However continued nucleation and growth of

intermetallic phases did not proceed through a surface-to-interior lithiation front mechanism described above for bulk Al materials (Fig. 2-3-5). Beyond the initial nucleation and growth events the continued formation of intermetallic phase instead propagated laterally across the anode surface. Additional nucleation events were only observed to occur at the boundary between lithiated and unlithiated material rather than continuing at previously isolated unreactive regions. The propagation of the lateral lithiation fronts was highly non-uniform with different sub-regions of the anode lithiating at very different rates. In contrast in-situ TEM of thin Si anodes in the same study revealed lithiation fronts that spread uniformly across the anode surface. The difference

in lithiation behavior between these two materials was again ascribed to the phase growth dependence in Al anodes on crystallization into β-LiAl. In comparison lithiation of Si

occurs through a solid-state amorphization mechanism as described previously in Ch. 2.2. A comparative schematic of the surface-to-interior lithiation front mechanism versus that of lateral lithiation front propagation for bulk versus thin film Al anodes is shown in Fig. 2-3-6.

Figure 2-3-6: Schematic of lithiation mechanisms and lithiated phase front progression (a) Surface-to-interior lithiated phase front progression characteristic of bulk Al anodes, (b) Lateral phase front propagation characteristic of nanostructured thin film Al anodes. Lines denote grain boundaries and half-circles denote nucleation points, from ref. 28.

Like Si and Sn based materials the focus on nanostructures has entirely driven the development to improve the performance of Al anodes because it is also a metal-alloying anode, and the volumetric expansion of 97% for LiAl formation is still relatively large compared to the small value of 6% for LiC6 formation in graphite anodes [3]. Therefore it is assumed that Al nanostructured anodes such as amorphous thin films, powders and nanowires will offer the same advantages relative to bulk Al anodes described previously for Si and Sn nanostructures, particularly in the accommodation of volume changes.

However in the literature nanostructured Al anodes have consistently shown high capacity initially with rapid capacity loss observed after a few cycles [14-24].

Currently there is significant debate in the literature concerning the dominant capacity loss mechanisms responsible for the continuing poor performance of Al nanostructured anodes. Firstly the volumetric expansion and contraction of intermetallic alloy formation and dissolution can impair stability of the solid-electrolyte interphase (SEI) layer present on the anode surface [29]. The SEI layer is formed initially from partial irreversible reduction and decomposition of the electrolyte. In graphitic anodic materials the volumetric expansion for lithiation is low at around 6% [3], resulting in minimal growth or change of the SEI layer beyond the first cycle. In Al like in other metal-alloying anodes such as Si the much larger volume changes can partially destroy the SEI layer present upon delithiation [29]. With continued expansion and contraction of cycling this will expose fresh Al material for continuous formation of thick SEI layers, causing significant permanent loss of lithium from the electrolyte. Secondly there is the pulverization of the active LiAl material [14, 19], which is considered by many to be the dominant failure mechanism in Al nanostructures because it is consistently observed in Li-Si and Li-Sn intermetallic alloy active materials. The progression of pulverization in Al anodes has been intensely studied through in-situ TEM of Al nanowires (NW) during repeated lithiation-delithiation cycles, shown in the series of images in Fig. 2-3-7 [19]. Fig. (a) shows a schematic illustration of the in-situ experimental setup. Fig (b) shows a pristine Al NW with diameter of 40 nm contacted with the Li2O/Li electrode to form an electrochemical device. Fig. (c) then shows the Al NW after the first lithiation with volume expansion observed in both radial and longitudinal directions. Progressing to the first delithiation stage (d) voids are formed indicated by red and yellow arrows. Continued delithiation in (e) enlarges these voids and forms new ones, indicated in blue arrows. After the second lithiation (f) the voids shrink and are partially healed. Switching back to the second delithiation in (g) increases the number and size of the voids. These trends continue until the NW is finally pulverized into nanoparticles (l, m), causing complete loss of electrical contact within the LiAl active material as well as to the supporting substrate.

Figure 2-3-7: Pulverization of a single aluminum NW upon electrochemical cycling, from ref. 19.

Another body of literature has emerged claiming that the dominant capacity loss mechanism is instead "lithium trapping", due to diffusion-limited Li transport in Al [24]. Like in Si anodes Al suffers from poor mass transport kinetics associated with the breaking and forming of chemical bonds upon electrochemical alloying of Li with Al. As described previously lithiation-delithiation of Al occurs through a mechanism of two LiAl phases (α and β) based on the lithium concentration available within the host

matrix. The reported diffusion coefficients of Li within α-LiAl and β-LiAl phases vary

significantly but are around 10-11 cm2/s and 10-9 cm2/s respectively [24]. These values are relatively slow compared to the diffusion coefficient of 10-7 cm2/s for Li in LiC6 in a graphite anode [1]. Practically speaking, the impaired Li transport in LiAl would limit the discharge rate and thus the power density. Three decades ago Owen et al. proposed a model to explain lithium trapping shown in Fig. 2-3-8 in terms of the charge and discharge processes [11]. For simplicity the volumetric changes of alloy formation and dissolution have not been included in the figure. Additionally it is assumed that the surface oxide layer plays no part in lithiation-delithiation and therefore merely acts as an ion conducting pathway to the Al core. The graphs in the bottom half show the Li concentration versus distance into the anode.

Figure 2-3-8: Owen's model for lithium trapping in charge/discharge processes of Al anodes, from ref. 11.

Initially the α-LiAl phase nucleates on the Al surface. As lithiation continues the α-LiAl

phase will progress into the Al bulk, while near the electrode surface the increasing Li concentration will convert α-LiAl phase to β-LiAl. Upon applying the opposite current

delithiation (oxidation) will begin from the LiAl/alumina interface, regenerating the Al from its surface and then progressing inwards. Therefore the recovery of lithium back into the electrolyte first happens at the electrode surface. Due to the differing diffusion coefficients of the two phases this gives rise to a lithium-rich β−LiAl layer trapped

between two lithium deficient α-LiAl layers. Additionally the lithium that is now trapped

can diffuse towards the surface or deeper into the Al bulk. Experimentally the phenomenon of lithium trapping has been shown by observing lithium present in Al electrodes even after the de-alloying step [25]. Considering diffusion is time dependent this capacity loss mechanism can become progressively worse with additional cycling. However recent evidence shows that it may be possible to circumvent this issue by minimizing the time during which the Al electrode is maintained at oxidizing potentials, because it allows less time for the trapped lithium to diffuse further into the Al bulk [24]. This is performed experimentally by cycling the Al electrode to a lower anodic (oxidizing) potential limit during delithiation.

Finally for Al anodes there is the point of contention regarding the influence of surface alumina layer on cycling performance and failure mechanisms. In other active electrode materials for lithium ion batteries such as LiCoO2, LiMn2O4 and MoO3 it is widely reported that the presence of an alumina layer can improve the durability and rate capability of the electrode during lithiation/delithiation [30-33]. However the underlying mechanisms of the alumina layer on these materials and Al anodes are well not understood in the literature. In Sn-based systems the surface oxide (SnO2) present natively on the anode is typically reduced first at a potential considerably higher before lithiation begins [4]. For Al anodes some studies assert that the oxide merely acts as an inert ion conducting pathway for lithium towards the Al core [24], as demonstrated above in the Owen's model. These studies show that changing the alumina thickness does not noticeably affect the electrochemical performance. Other studies show that the alumina is first irreversibly lithiated into a super-hard Li-O-Al layer [19], similar to lithiation of

SiO2 that occurs in Si anodes [34]. This process occurs first prior to lithiation of the Al core. The resulting Li-O-Al layer has poor electrical conductivity compared to the original alumina layer, and is therefore considered a source of permanent capacity loss. However the Li-O-Al layer acts to contain the pulverization of electroactive materials, by enlarging through elastic and plastic deformation to act as a solid electrolyte with exceptional mechanical robustness and ion conduction. Shown in Fig. 2-3-9 are in-situ TEM images and EELS maps of a pristine Al nanowire anode with a native 4 nm oxide layer that was lithiated into a thicker 5 nm Li-O-Al layer [19].

Figure 2-3-9: Evolution of the surface Al2O3 layer to the Li-O-Al layer. (a) Pristine

Al nanowire with 4 nm native Al2O3 layer, (b) Lithiation of the surface layer, whose

thickness was increased to 5 nm, (c) The Al nanowire with the lithiated surface layer, (d-f) EELS maps of Li, Al, O respectively, indicating that the surface Al2O3

layer had evolved to Li-O-Al after lithiation, from ref. 19.

The continuing poor performance for lithiation-delithiation of Al nanostructured anodes suggests that the alternative of bulk Al anodes may be preferable. Bulk Al materials offer advantages of ease of preparation, mechanical robustness and improved conductivity over nanostructures. For this purpose it is important to consider the processing methods used

to manufacture commercial bulk aluminum. As described in Ch. 2.3.1 there are a wide variety of processing techniques available to modify the mechanical properties of bulk Al metal through mechanical, thermal or alloying methods. Therefore commercial bulk Al metal can offer a wide variety of mechanical properties.

As described previously the mechanism of lithiation-delithiation in metal-alloying anode materials such as Al depends on the density and mobility of dislocations within the host [25-27]. One would expect the processing methods on bulk Al materials will determine the arrangement and mobility of dislocations within the host [5-7], which will affect the relative ease of plastic deformation of the resulting LiAl alloy. Therefore one should expect that the resulting mechanical properties of bulk Al materials should influence these various lattice contractions and expansions for nucleation and growth of intermetallic phases. As an anode in lithium-ion batteries these properties will in turn affect the material's response to the volume changes of lithiation-delithiation, and ultimately electrode degradation and failure [35]. Additionally these mechanical properties should affect the properties of any surface oxide present on the Al material. The influence of mechanical properties and the surface oxide of Al materials on lithiation-delithiation, specifically capacity loss mechanisms such as solid-electrolyte interphase (SEI) formation, surface oxide lithiation/reduction, pulverization of active materials and mass transport limitations of Li in the bulk material (lithium trapping) are key issues that will be investigated in this thesis work beginning in Ch. 4.1.