Interfacial fracture mechanisms in solid solution directionally solidified eutectic oxide composites

Interfacial fracture mechanisms in solid solution directionally solidified eutectic oxide composites

Acta Materialia 52 (2004) 3781–3791 www.actamat-journals.com Interfacial fracture mechanisms in solid solution directionally solidified eutectic oxide...

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Acta Materialia 52 (2004) 3781–3791 www.actamat-journals.com

Interfacial fracture mechanisms in solid solution directionally solidified eutectic oxide composites L.N. Brewer 1, M.U. Guruz 2, V.P. Dravid

*

Department of Materials Science and Engineering, Northwestern University, Evanston, IL 60208, USA Received 19 December 2003; received in revised form 19 April 2004; accepted 20 April 2004 Available online 17 June 2004

Abstract The interfacial fracture behavior of the solid solution directionally solidified eutectic oxide, Co1  x Nix O/ZrO2 (CaO), is investigated via indentation testing. An abrupt transition from interfacial delamination to interfacial penetration is observed as a function of NiO fraction (x > 0:33). The use of a focused ion beam technique is introduced as a means for exploring sub-surface cracking in brittle materials. The sub-surface observations revealed a compositional transition from energy dissipative mechanisms (e.g., delamination, secondary cracking) for cracking to brittle cracking behavior, in the Co1  x Nix O phase. It is proposed that the transition in interfacial fracture behavior for the directionally solidified eutectic material is the result of competing dislocation-based crack nucleation mechanisms. The observations and analyses have significant implications for fracture behavior in composites with one phase exhibiting (pseudo-) plastic behavior. Ó 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Interfaces; Eutectic; Fracture; Indentation; Focused ion beam

1. Introduction Directionally solidified eutectic (DSE) oxides have been previously investigated as ultra-high temperature (>1500 °C) composites and as model systems for fundamental materials science studies [1,2]. Materials such as YAG/Al2 O3 (YAE) [3–8], Al2 O3 /ZrO2 (AZE) [9,10], and GdAlO3 /Al2 O3 [11] have shown promising creep resistance and high temperature strength, both in fiber form and as monolithic materials. However, it is the low fracture-toughness of these materials (4–5 MPa m1=2 ) at low temperatures that has hindered their development as reliable structural materials [12]. It is primarily the inability of the internal interfaces in oxide–oxide DSE materials to deflect cracks or to delaminate interfaces that is blamed for the poor fracture toughness of the composite as a whole. However, the exact nature of interfacial fracture has not been com* Corresponding author. Tel.: +1-847-467-1363; fax: +1-847-4917820. 1 Now at Sandia National Laboratory, Albuquerque, New Mexico. 2 Now at Hitachi Hard Disk, San Jose, California.

pletely determined for either engineering or model DSE oxides. The central interfacial fracture behavior of interest in DSEs is the competition between the penetration of the interface by a nucleating or impinging crack versus the deflection/delamination of this crack along the interface (Fig. 1) as has been discussed previously by several authors [13–16]. Indentation fracture studies on the engineering oxide DSE’s YAG/alumina (YAE) and alumina/zirconia (AZE) have indentified interfacial penetration as the dominant interfacial fracture mode [10,12,17]. Interfacial penetration is the dominant observed mode in YAE fibers, which also exhibit brittle fracture surfaces [12]. Some interfacial crack deflection was observed by enriching the YAG/alumina interfaces with rare-earth cations such as Pr and Ce, but an equal number of interfacial penetration events were also observed [17]. Studies on AZE with partially stabilized ZrO2 (Y2O35.3 mol%) claimed to observe interfacial penetration (primarily for indents on transverse sections) and some interfacial deflection (primarily for indents on longitudinal sections) [10]. In summary, indentation studies on interfacial fracture in two of the major engineering oxide

1359-6454/$30.00 Ó 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2004.04.016

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Interfacial Delamination

Impinging Crack

Interfacial Penetration

Fig. 1. Schematic illustrating competition between interfacial penetration versus deflection/delamination of a crack impinging on an interface between dissimilar solids after He and Hutchinson (1989) [13].

ZrO2 (CaO) system. We use a combination of traditional microindentation techniques in combination with novel sub-surface observations of cracks using a focused ion beam (FIB) technique. In addition, we investigate the compositional dependence of hardness as a measure of room temperature plasticity in Co1 x Nix O single crystals and the Co1 x Nix O/ZrO2 (CaO). Both Vickers indentation and nanoindentation are used to measure microstructural-average plasticity (tens of microns) and spatially resolved plasticity across single interfaces (submicron). The goal is to develop a phenomenological model for understanding the interfacial fracture problem in this DSE material.

2. Experimental methods 2.1. Single crystal and DSE growth

DSEs have shown a dominance of interfacial penetration by indentation-induced cracks with little to no deflection of cracks by interfaces. A limited number of studies have also been performed on the interfacial fracture behavior in model oxide DSE systems. The extensive study by Hulse and Batt [1] in the early 1970s surveyed the behavior of a number of oxide–oxide DSEs, most of which demonstrated interfacial penetration by cracks. The notable exception was the PbO–3PbO Nb2 O5 , which exhibited extensive interfacial delamination around Vickers indentations, although this material was not pursued further due to its low melting point, lack of chemical stability, and poor transverse fracture response [1]. Indentation fracture studies on NiO/Y2 O3 , NiO/ ZrO2 (CaO), and NiO/CaO have demonstrated interfacial penetration as the dominant mechanism in the material [18–20]. These observations were again attributed to strong interfacial adhesion with the conclusion that the interfacial adhesion was perhaps greater than the cohesive energy of either phase [19]. The solid solution DSE oxide, Co1 x Nix O/ ZrO2 (CaO), is a novel material which has been recently developed for the study of interfacial fracture structure– property relationships in oxide composites [21]. This DSE has an identical microstructure to NiO/ZrO2 (CaO) and also has an interfacial stacking sequence involving the {1 1 1} planes of Co1 x Nix O and the {1 0 0} planes of ZrO2 (CaO) as in NiO/ZrO2 (CaO) [22]. These structural similarities beg the question of whether or not the interfacial structure of the solid solution DSEs will produce similar interfacial fracture behavior as that observed in NiO/ZrO2 (CaO). A systematic change in interfacial fracture with composition might provide a method for tailoring mechanical behavior in DSE oxide composites. This paper investigates the mechanisms responsible for interfacial fracture behavior in the Co1 x Nix O/

Samples of Co1 x Nix O/ZrO2 (CaO-15 mol %) were directionally solidified using the optical floating zone method at the Laboratoire de Physico-Chimie de l’Etat Solide, Universite de Paris-Sud (Orsay, France) as has been described in more detail elsewhere [21]. Appropriate fractions of CoO, NiO, ZrO2 , and CaCO3 were combined into a powder mixture, calcined, isostatically pressed, and finally sintered. The sintered rods were directionally solidified using an NEC commercial double ellipsoidal image furnace. Controlled atmospheres of Ar, CO2 , and air were used to suppress the formation of Co3 O4 during growth. Single crystals of Co1 x Nix O were grown under the same conditions using a CSI (FZT-4000-H-II-PP) furnace that has been described in the related paper by Brewer et al. [23]. 2.2. Vickers indentations Samples were prepared by cutting transverse sections from single crystal and DSE rods using a low-speed diamond saw. These sections were polished to a 1 lm finish and then finished with 0.05 lm SiO2 suspension. The surface orientation of the single crystal samples was confirmed to be h1 0 0i using a combination of electron back-scattered diffraction (EBSD) and Laue X-ray diffraction. Vickers indentations were made using a commercial, diamond tipped microindenter (Beuhler, MicrometÒ II Microhardness Tester, Model # 1600– 9000). Loads from 50 g (0.5 N) to 1 kg (10 N) were applied with hold times of 5 and 10 s. Indents were made with the diagonals of the Vickers indenter oriented along the h1 0 0i surface axes for the single crystal samples. The indenter diagonals were aligned along the lamellae for the DSE samples. Optical and SEM images were taken of the indents to observe the shape of the indent and the presence of any cracks immediately following the indentation experiment.

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The sub-surface cracking behavior was observed by making trenches 20 lm in width by 10 lm in length by 20 lm in depth both parallel and perpendicular to the lamellae in indented material using the Hitachi FB2001A FIB (Fig. 2). This FIB uses a liquid Ga ion source with an accelerating voltage of 30 kV. The samples used for this experiment were in the as-indented condition. The sub-surface cracking behavior was observed by placing the ion-milled samples in the Hitachi S4500 cFEG SEM and tilting the sample to 50°, using secondary electrons as the imaging signal. Because of interest in the transmission of cracks from one layer to the next, a sequential milling procedure was used (see Fig. 2 for reference) consisting of: (1) milling an initial trench, (2) observation of cracking in an individual layer ‘‘A’’, (3) subsequent milling of layer ‘‘A’’, (4) observation of sub-surface cracking in the adjacent layer ‘‘B’’, and so forth. 2.3. Nanoindentation Nanoindentations were made with a commercially available nanoindenter (Hysitron Inc. TriboscopeÒ System, Item 5-099 combined with Digital Instruments NanoscopeÒ II base) using a diamond Berkovitch tip calibrated on a fused silica standard. Samples were prepared in the same fashion as for the Vickers indentation measurements. Loads of 500, 1000, 2000, and 3000 lN were applied with at least 30 indentations performed per composition. Series of nanoindentations across interfaces were spaced 200 nm apart using a load of 500 lN. The resulting indentation curves were analyzed using the Oliver–Pharr method on the unloading portion of the curve. [24] The use of the NanoscopeÒ II AFM base in this experimental setup allowed the in situ imaging of the nanoindentations for verification of their placement within the microstructure.

Fig. 2. Trenches milled by FIB at a Vickers indent in Co0:9 Ni0:1 O/ ZrO2 (CaO).

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In addition to the hardness, the plasticity factor (PF) was calculated as defined by Pharr [24],

P:F ! 0;

hfinal ; hmaximum Perfectly Elastic Deformation;

P:F ! 1;

Perfectly Plastic Deformation;

Plasticity Factor ¼

ð1Þ

where hmaximum is the indentation depth at maximum load and hfinal is the indentation depth after the load is removed. The quantity is readily measured from the indentation data and provides more information about the nature of the deformation than the hardness number alone.

3. Results 3.1. Indentation fracture When loaded under identical conditions, the indentation fracture patterns for NiO/ZrO2 (CaO) and CoO/ ZrO2 (CaO) (Figs. 3 and 4) are distinctly different.

Fig. 3. Interfacial fracture behavior for NiO/ZrO2 (CaO). (a) Optical image of whole indentation. (b) SEM image with bright phase as ZrO2 (CaO) and dark phase as NiO. Note that the crack penetrates every interface for several layers.

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extended cracks, particularly for the indentation diagonal most closely parallel with the lamellae (Fig. 4(a)). However, extensive interfacial delamination was observed for the indentation diagonals most closely perpendicular to the lamellae. At higher magnifications, the interfacial delamination in CoO/ZrO2 (CaO) is revealed to be fairly complex. (Fig. 4(b)). The crack system is composed of interfacial delamination at every other interface. The delamination is only observed at the interface going from the CoO phase to the ZrO2 (CaO) phase in the direction moving away from the indent (i.e., a Co1 x Nix O ! ZrO2 (CaO) interfacial fracture event). The interface proceeding from ZrO2 (CaO) to CoO in this direction either penetrates a crack or remains well bonded. In addition, a set of short cracks, perpendicular to the lamellae, nucleate almost exclusively in the ZrO2 (CaO) phase. It is not clear if these ‘‘secondary’’ cracks are indeed secondary in formation to the delamination or if they are generated simultaneously or even precede the delamination events. The interfacial fracture response of the solid solution eutectics showed a distinct transition as a function of composition. For mole fractions of NiO less than 20%, the interfaces perpendicular to the corner of the Vicker’s indent delaminated (Fig. 5(a)). However, for interfaces at shallow angles to the indenter corner, interfacial penetration and extended cracks were observed. For NiO fractions of greater than 20%, cracks penetrated interfaces with little observable deflection (Fig. 5(b)). 3.2. Sub-surface observations

Fig. 4. Interfacial delamination and secondary cracking in CoO/ ZrO2 (CaO). (a) Optical image of whole indentation. Note that delamination takes place only at every other interface. (b) SEM image with bright phase as ZrO2 (CaO) and dark phase as CoO. Note extensive delamination parallel to lamellae in CoO/ZrO2 (CaO) (SEM back-scattered electron images).

Vickers indents on NiO/ZrO2 (CaO) create cracks which emanate from the indent to a length equal to or greater than the width of the indent (Fig. 3(a)). The length of these cracks is independent of the orientation of the indent with respect to the microstructure. The key feature of the extended cracks in NiO/ZrO2 (CaO) is the pronounced interfacial penetration (Fig. 3(b)). The crack propagates through one interface after another, seemingly unperturbed. Vickers indents on CoO/ ZrO2 (CaO) under identical conditions produce some

The sub-surface cracking of NiO/ZrO2 (CaO) revealed a much more complex behavior than was observed from the surface. In Fig. 6(a), the line trace of a crack plane’s intersection with the surface is visible. With this information alone, one might conclude that this crack is of radial-median type because of the way the crack extends away from the center of the indentation. However, looking subsurface at the side of a zirconia layer revealed that the crack is in fact more sublateral in nature (parallel to the surface of the sample). Upon removal of three additional lamellae and inspection of the side of a NiO layer (Fig. 6(b)), an extension of the same crack system was observed. Again, this crack demonstrates a sublateral character below the surface that might be mistaken as radial median from surface observations only. The crack front is also now further below the surface of the sample, suggesting that the crack front is propagating deeper into the sample away from the indentation. Another important observation here is that ‘‘the same’’ crack front is observed in the NiO layer as in the zirconia layer, indicating a continuous crack front from one layer to the next, and therefore interfacial penetration as the dominant interfacial fracture mode.

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Fig. 5. Transition in interfacial fracture behavior for the SS-DSE, Co1 x Nix O/ZrO2 (CaO), with NiO fraction, x. Note delamination along lamellar interfaces in Co0:9 Ni0:1 O/ZrO2 (CaO) (optical image with dark phase as ZrO2 (CaO) and bright phase as Co0:9 Ni0:1 ).

Fig. 6. Subsurface observations of indentation cracks in NiO/ZrO2 (CaO). Note that the same crack system is observed in both phases (SEM image).

Subsurface observations on CoO-rich compositions reveal distinctly different subsurface cracking behavior. Fig. 7(a) displays the side of a zirconia layer at a trench made parallel to the lamellae, the direction for which delamination is observed. Surface observations show cracks in zirconia running perpendicular to the lamellae and parallel to the surface normal. However, a far more complex system of cracks is observed beneath surface than the line traces observed on the surface. Not only is the crack system in the zirconia layer highly branched, but it also appears from this image that the cracks are nucleated below the surface and propagate outwards and upwards, eventually intersecting the surface of the sample. Removal of this zirconia layer with the FIB and observation of the adjacent Co0:9 Ni0:1 O layer reveals a further departure from the fracture behavior observed for NiO/ZrO2 (CaO) (Fig. 7(b)). On the surface, cracks were observed in the zirconia layer, not in the Co0:9 Ni0:1 O layer, but then again in the next zirconia layer. Subsurface inspection of the Co0:9 Ni0:1 O layer in Fig. 7(b) revealed no cracks in contrast to the observation in the NiO layer. The removal of this Co0:9 Ni0:1 O layer and inspection of the next zirconia layer (Fig. 7(c))

again displayed the presence of similar surface and subsurface cracks as observed in Fig. 7(a). A final layer removal reaffirmed the absence of any cracks in the Co0:9 Ni0:1 O layer (Fig. 7(d)). 3.3. Indentation hardness 3.3.1. Vickers hardness The Vickers hardness of the solid solution single crystals, Co1 x Nix O, increased with x by more than a factor of two (Fig. 8). The measured hardness values for CoO and NiO are in agreement with previous Vickers hardness measurements on single crystals of these materials, 2.97 and 5.3 GPa, respectively [25]. Steep increases in hardness were observed at either end of the solution, exhibiting solid-solution hardening like behavior. The evolution of hardness in the middle of the solid solution followed a rule-of-mixtures type of behavior. These trends were observed for all levels of load, although the actual hardness values increased for lower loads and decreased at higher loads due to the onset of cracking, as been previously observed for Vickers indentation experiments [26].

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Fig. 7. Layer-by-layer sequence in FIB-indent Co0:9 Ni0:1 O/ZrO2 (CaO) sample. Black arrow represents equivalent positions in each micrograph. Higher contrast version included for each image below the primary image.

700

7

600

6

500

5

400

4

300

3

200

Hardness (GPa)

Hardness (Vicker's-200gf)

Hardness across the NiO-CoO Solid Solution

2 0

0.2

0.4 0.6 0.8 Mole Fraction NiO

1

Fig. 8. Vickers hardness measurements Co1 x Nix O single crystals. Error bars represent 2 SD from the mean value.

The Vickers hardness of the SS-DSE materials also increased with increasing NiO fraction, but with greater scatter than observed in the Co1 x Nix O single crystals (Fig. 9). The hardness values are substantially harder (5– 9 GPa) than for the single crystals alone (3–6 GPa), as is expected due to the higher hardness of cubic zirconia (12–15 GPa) phase. The hardness of zirconia should not change with NiO fraction in the SS-DSE nor does the volume fraction of zirconia increase substantially with composition; thus, the increase in hardness with composition is most likely due to the increase in hardness from the Co1 x Nix O phase, as is observed for the single

Vickers Hardness vs Composition for SS-DSE's 9 8 7 6 5 0

0.2

0.4

0.6

0.8

15 10 5

1

0

crystals. No measurable difference was observed in hardness measurements made on transverse versus longitudinal DSE sections. 3.3.2. Nanoindentation hardness and plasticity factor The hardness values of the Co1 x Nix O solid solution increase with increasing NiO fraction (Fig. 10). The dependence of hardness upon composition is nearly linear, as opposed to the nonlinear dependence observed for the Vickers indentations. The nanoindentation hardness values are approximately double the value of the Vickers hardness values. Furthermore the hardness for the nanoindentations decreased with increasing load from the minimum nanoindentation load of 500 lN to the maximum of 3000 lN, approaching the hardness values from the Vickers measurements. These increases in hardness with lower loads are in keeping with the indentation size effects observed in metallic thin films by Nix and co-workers [27,28] and recently for the isostructural alkalki halide crystals, LiF and KCl [29]. The nanoindentation hardness values for the SS-DSE materials exhibited slight shifts from the hardness values measured for the single crystals (Fig. 11). The nanoindentation hardness values for the Co1 x Nix O phase were larger (greater than one standard deviation) in the

Plasticity Factor

Hardness (GPa)

8 6 4 2 0.6

0.6

0.8

1

DSE material (8.5–12 GPa) than in the single crystals (5–10 GPa) with the hardness increasing somewhat linearly with increasing NiO fraction. The ZrO2 (CaO) hardness values ranged from 20 to 30 GPa, much harder than the Co1 x Nix O phase for all compositions, but in keeping with measurements in the zirconia literature for small indentations [30]. The PF exhibited a strong dependence upon composition across the Co1 x Nix O solid solution single crystals (Fig. 12). The PF value for CoO was approximately 0.7, which indicates that the degree of plastic deformation is large enough that significant pile-up of material takes place about the indenter [24]. This level of pile up is normally observed for metals such as Al. The PF decreased sharply with the addition of NiO to the solution. The PF for the individual layers in the SS-DSE material reduced with increasing NiO fraction in a similar manner as the single crystals of the Co1 x Nix O phase. The PF for the ZrO2 (CaO) phase was in the range from 0 to 0.2, with no discernable dependence upon composition along the solid solution. The hardness (Fig. 13(a)) and PF (Fig. 13(b)) can be plotted as function of position across the interface (Fig. 13(c)). For the representative composition of

0.8

Plasticity Factorvs Composition

0.7

10

0.4

0.4

Fig. 11. Comparison of nanoindentation hardness values for single crystals and DSEs.

Vicker's Hardness (GPa) -2N Load Nanoindentation Hardness (GPa)-0.001N Load Vickers Literature Values

0.2

0.2

Mole Fraction NiO

Fig. 9. Vickers hardness measurements on transverse sections Co1 x Nix O/ZrO2 (CaO).

0

Single Crystal DSE

20

Fraction NiO

12

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Hardness vs Composition

25

10

Hardness (GPa)

Vickers Hardness(GPa)-200gf

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1

Mole Fraction NiO Fig. 10. Comparison of Vickers and nanoindentation hardness measurements on Co1 x Nix O single crystals.

Single Crystal DSE

0.6 0.5 0.4 0.3 0.2 0

0.2

0.4

0.6

0.8

1

Mole Fraction NiO Fig. 12. Plasticity factor versus composition for single crystals and DSEs.

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Fig. 13. Changes in plastic response across interfaces in the DSE Co0:666 Ni0: 333 O/ZrO2 (CaO): (a) hardness; (b) plasticity factor; (c) nanoindentation points across interfaces in a solid solution directionally solidified eutectic. The large indents at either end were used for marking purposes.

Co0:66 Ni0:33 O/ZrO2 (CaO), one observes an abrupt (<200 nm) increase in hardness and decrease in PF from the Co0:666 Ni0:333 O phase to the ZrO2 (CaO) phase. Within at least 200 nm, there does not appear to be a measurable interfacial influence on the mechanical response within a given layer.

4. Discussion The compositional transition in interfacial fracture behavior observed in this study might be explained by compositional shifts in interfacial bonding, residual stress levels, or room temperature plasticity, but it appears that room temperature plasticity is the dominant factor. No data currently exist for the change in interfacial adhesion energy with composition for the

Co1 x Nix O/ZrO2 (CaO) interface. However, eutectic interfaces are by definition stable, low energy interfaces, which are strongly bound. This strong interfacial bonding is indirectly observed for all compositions by the ability of cracks at shallow angles to penetrate across the interfaces unperturbed. A large change in interfacial bonding with composition should exhibit a propensity for interfacial deflection at one end of the solid solution. Lastly, large residual stresses (see below) have been observed in many ZrO2-based DSEs, indicating effective load transfer and the ability of the interfaces to withstand such large residual stresses. This further provides indirect evidence for well-bonded interfaces. Large (1 GPa) residual stresses have been observed in NiO/ZrO2 (CaO) [31] and more recently in the compositions Co1 x Nix O/ZrO2 (CaO) x ¼ 0, 0.333, and 0.5 [32]. In each case the sign of the residual stresses acting on the interfacial plane is the same: tensile for Co1 x Nix O and compressive for ZrO2 (CaO). Although the residual stress levels have not been measured for CoO-rich compositions, the coefficients of thermal expansion [23] and calculated residual stress levels [32] predict the same signs and magnitudes of residual stresses all the way across the solid solution. If these predictions are accurate, then residual stresses would not be the dominant controlling factor in the interfacial fracture transition as there would need to be a change in the sign of the residual stress to explain the transition (stress in Co1 x Nix O going from tensile to compressive with increasing CoO). A compositional change in the room temperature plasticity of the Co1 x Nix O phase is the most plausible reason for the observed transition in interfacial fracture. Oxides such as MgO [33], NiO [34], and CoO [25,35] and metal-halides such as NaCl and LiF have demonstrated strains at failure of up to 4% at room temperature under compression [25,34]. The majority of these experiments has been done on single crystals oriented with h1 0 0i as the compression directions. Related studies have focused on the differences in compression behavior for crystals oriented along the h1 1 1i direction in NiO and CoO [34,36]. TEM studies of dislocation activity after compression have revealed the primary slip systems to be {1 1 0} h1 1 0i for rocksalt structures at room temperature [36,37]. These studies have shown evidence for activation of the {1 0 0} h1 1 0i the slip system at room temperature for CoO [36], but activation of this slip system has only been observed at higher temperatures (500 °C) for NiO and MgO [34]. This difference in active slip systems for CoO and NiO may explain the large increase in hardness and decrease in PF between the two ends of the solid solution. The lack of sub-surface cracking in the Co1 x Nix O layers also supports the idea that the CoO-rich layers deform under indentation by plastic flow, while the NiO-rich layers deform by brittle fracture.

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To construct a hypothesis around plasticity-based interfacial fracture, we suggest dislocation-based crack nucleation mechanisms as the fundamental difference between the interfacial fracture modes observed. We are focused on the nucleation of cracks in solid solution DSE oxides, and in particular on the Co1 x Nix O ! ZrO2 (CaO) interfacial fracture event. Two separate models need to be considered for the nucleation of cracks in elastoplastic materials: the Zener–Stroh–Koehler (ZSK) crack [38,39] and the Cottrell–Keh (CK) crack [40,41]. The ZSK crack is used to explain the nucleation of cracks in metals by pile-up of dislocations at grain boundaries or precipitate boundaries. The ZSK crack nucleates when the tensile stress field generated by dislocation pileup at the grain boundary exceeds the Griffith criterion. In 1958, Cottrell proposed the nucleation of a crack between intersecting slip planes in BCC metals as a means for explaining brittle, transgranular fracture [40]. A similar sort of crack was introduced by Keh et al. [41] that incorporates an energetically more favorable reaction for dislocation coalescence in rocksalt structures, a  a a ½0 1 1 þ ½1 0  1 ! ½1  1 0; 2 2 2

ð2Þ

where a=2h1 1 0i Burgers vectors are along {1 1 0} planes. This type of crack (called here the Cottrell–Keh or CK crack) has been observed as the dominant crack nucleation event in LiF, MgO, and NiO single crystals during indentation of the {1 0 0} surface [20,41]. ZSK cracks will only be observed in polycrystals or multiphase materials as they require an interface on which to nucleate, while the CK crack readily forms in single crystals needing only dislocation locking between slip planes to nucleate. In general, the CK crack requires fewer dislocations to form than the ZSK crack. This statement is particularly true for materials with a limited number of slip systems because they run the greatest risk of experiencing significant dislocation pileup from intersecting slip planes. In a study of the indentation fracture of the NiO/CaO DSE, both kinds of cracks were observed in equal numbers, thus providing no insight to the competition between interfacial fracture mechanisms [20]. Based on these observations, we propose that the transition in interfacial fracture behavior in the Co1 x Nix O/ZrO2 (CaO) system is well described by the competition between ZSK and CK cracks (Fig. 14). Note first that the delamination/deflection event always takes place as the deformation front proceeds from the Co1 x Nix O phase to the ZrO2 (CaO) phase, from a elastoplastic phase to a brittle elastic phase. The CoO-rich compositions could permit propagation of dislocations along readily available slip planes towards the Co1 x Nix O/ZrO2 (CaO) interface. The resulting dislocation pile-up at the interface is sufficient for the nu-

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Cottrell Crack

Elastic Solid (Zirconia)

ZSK Crack Elastoplastic Solid CoNiO Fig. 14. Schematic illustration of dislocation pile-ups resulting in microcrack nucleation at plastic–elastic interface (ZSK crack) and at glide plane intersections (Cottrell–Keh crack).

cleation of a ZSK crack. For NiO-rich compositions, the lack of available slip planes could cause a CK crack to nucleate prior to a nucleation event for the ZSK crack. The transition in interfacial fracture is the compositional point at which dislocation coalescence is favored at interlocking slip planes rather than at the interface. This hypothesis will be further investigated by direct observations of the dislocation interactions near cracks and near the interfaces in these solid solution DSE oxide materials. Even though the present observations are limited to oxide DSEs, the analysis presented should be applicable to general brittle–ductile composite systems with strongly bonded interfaces. In such systems, one would expect to see energy dissipative mechanisms, such as interface delamination and secondary cracking. Further, we believe that the interfacial delamination in such system would be at alternate interfaces, i.e., when the crack propagates from ductile to brittle phase. Lastly, this work also hints at the potential to modify fracture behavior by appropriately doping one of the phases in composites to render it more plastic.

5. Conclusion This paper has investigated the interfacial fracture behavior of the Co1 x Nix O/ZrO2 (CaO) SS-DSE. Using indentation fracture measurements, a distinct transition from interfacial delamination to interfacial penetration was observed with increasing NiO fraction. An important observation from these experiments was that only the Co1 x Nix O ! ZrO2 (CaO) interface delaminated and not the ZrO2 (CaO) ! Co1 x Nix O interface. Subsurface observations using a novel FIB technique demonstrated the absence of cracking in CoO-rich layers and the presence of cracking in NiO-rich layers. These observations suggest plasticity of the Co1 x Nix O layer

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as the controlling mechanism for the interfacial fracture behavior. Nanoindentation results support a systematic change in Co1 x Nix O layer plasticity as evidenced by increasing hardness and decreasing plasticity factor with increasing NiO fraction. The hypothesis of competing dislocation-based crack nucleation mechanisms has been introduced as a means of explaining the observations in this system. This hypothesis contends that CoOrich layers potentially posses sufficient plasticity at room temperature to nucleate an interfacial pile-up crack.

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