Prospect of Graphene Interface Control for Tougher Ceramic Composites

Graphene-Ceramic Composites (GCCs) have been little studied compared to graphene-polymer composites [1]. Recent reviews [2,3] indicate that both mechanical and electrical property ceramic improvements can be obtained by mixing small quantities, typically 1 to 15% of graphene material with a ceramic precursor, then compacting and sintering. The greatest effect is on the electrical properties. The electrical conductivity of a material was first shown to rise by several orders of magnitude for only a 1% volume addition of graphene as in polymer composites [4] but the stiffness, strength and toughness only increased by 20-160% or so at 5% addition, a rather minor improvement compared to significant increases caused by slight ceramic process changes. Some crack bridging and pull-out mechanism was observed by electron microscopy in graphene-alumina composites, though the effects were modest. Surface friction and wear improvements of around 100% were also notable. This paper seeks to show that much higher toughness increases might be produced using the method pioneered by Clegg et al [5], where the graphite interlayers are replaced with graphene to produce improved ordered interfaces with reliable coverage and consistent interface fracture energy, enabling an increase in the fracture resistance of the ceramic by two orders of magnitude. DOI : Coming Soon Corresponding author: Kevin Kendall, University of Birmingham, B15 2TT, UK Email: kevin.kendall@adelan.co.uk

less than 90% glass because of the polymer interlayers, so is not truly ceramic. The concept outlined here is one where the polymer interlayers in bullet-proof glass are replaced by controlled ceramic interfaces, made with nanoparticles like graphene. Previous work has shown that it is essential that any such film should be continuous with no gaps to allow crack escape into the underlying ceramic.
By toughness, we mean an increase in energy dissipation, that is an increase in the resistance to fracture, R, when cracks attempt to run catastrophically through the material. Energy dissipation slows down the crack, producing stable or graceful collapse as opposed to dangerous, explosive shattering. The fibre toughening effect occurs because the crack is made to cross many interfaces between fibres and matrix, with certain types of interface giving substantial energy loss. This principle has been known since the invention of When such materials were accidentally baked in an inert atmosphere in 1958, the resin matrix pyrolysed to carbon to produce a carbon-carbon composite with interesting strength and toughness properties [6,7]. But the difference between polymer-composites and ceramic composites is the large sintering shrinkage during ceramic manufacture. The reinforcing fibres cannot accommodate such shrinkage strains. Now, newly developed matrix processes, including shrink-free chemical vapor deposition of the carbon, give more high-quality carbon-carbon composites for technical applications [8] but at very much greater processing costs.
Transformation toughening is exemplified by the performance of partially stabilised zirconia eg 3YSZ (3%mol Yttria Stabilised Zirconia) described by Garvie and colleagues in the early 70s [9,10]. Analysis of cracks moving through this metastable material showed that energy was dissipated at the crack tip due to the change in structure and 3% volume expansion of the partially stabilised zirconia crystals under the intense crack tip conditions. An order of magnitude increase of fracture energy was possible from this mechanism.
In the invented process for tough ceramics dating back to the 90s [5,11] In other words, the condition for a Griffith crack to be deflected at an adhesive interface is that the adhesive interface fracture energy must be about ten times less than the ceramic fracture energy. Interface strength cannot enter this theory.
These ideas are relevant to laminar interfaces in composites because they were used successfully in [5,11]    showing increase of force with speed and higher force for delamination crack than for peeling. The lines are equations 1) and 2) with R replacing W.  [18][19][20].
When a smooth surface revealed by delamination heals in a new position, more energy must be dissipated to break the interface repeatedly, giving increased energy loss and therefore more toughness [13], as indicated by phenomena observed in fracture of nacre, which has a laminated structure with low interface fracture energy.

Graphene Interfaces
The major difficulty in this area is that very little is known about interfaces between ceramics and graphene and this is an area where significant further study is required. Measurements of the interface adhesion fracture energy, its dependencies on crack speed and temperature, the influence of the atoms attached to the graphene platelets, the effects of processing temperatures and environments; none of these are known or understood at present.

One exceptional paper appeared in 2010 from
Wan's group [21]. A thesis from Li, one of his PhD students, puts the ideas in timely perspective [38].
50nm diameter gold nanoparticles were spread thinly on a (100) silicon wafer oxidised with a layer 280nm thick.
Graphene sheets were mechanically cleaved from the surface of highly oriented pyrolytic graphite (HOPG-ZYH grade from NT-MDT Co., Moscow, Russia) using the celebrated method of Scotch tape peeling [22,23]. The tape carrying the adhered small flakes of graphite was then brought into contact with the Si substrate in dry condition, trapping a number of gold particles at the graphene-silicon interface. The thick silicon substrate suffered negligible deformation and was taken to be rigid, while the thin graphene sheet bent and formed a blister, in which the van der Waals forces pulling the graphene flake into contact with the wafer were balanced at equilibrium by the elastic reaction of the deformed flake. The measured interfacial adhesion energy R was calculated from the particle height and blister radius measured by scanning electron microscopy using the thin clamped membrane equation and energy balance crack equation [24] R = Eh (w/r) 4 /16 (4) where E=0.5TPa was the elastic modulus of the graphene sheet, h=1.7nm its thickness, 5 layers, w the diameter of the gold particle and r the blister radius. R was found to be 0.151±0.028 Jm-2 .
It is difficult to say how near this value of R=0.151 Jm -2 is to the equilibrium work of adhesion W.
To understand that, and to see the effects of crack speed, it would be necessary to observe the peeling and healing of the graphene sheet on the wafer with time and environmental conditions, as was achieved with elastomeric materials many years ago [25]. Of course, placing a graphite flake on a substrate under ordinary atmospheric and rough conditions is unlikely to give a realistic value of R = W, due to the presence on both surfaces of adsorbed chemical groups (Fig 4).
Since 2010, several interesting papers have followed to consider the interfacial fracture energies of graphene in contact with oxide and other surfaces.
Reference [26] was the first report of interface fracture energy for a large surface area graphene monolayer grown by CVD on copper. One interesting development was to consider a more standard interface fracture energy test than the blister method, which is not generally applicable to laminate systems. The double cantilever beam (DCB) method in contrast is widely favoured for laminates and has a long track-record for polymer composites [27]. A target substrate was bonded with epoxy resin to the graphene surface ( Fig 5) and bending cleavage forces applied. The value of R=0.72±.07 Jm -2 was determined under ambient room conditions. However, the velocity of the crack was not determined, and the dependence of R on speed and temperature is unknown. If epoxy resin had been used to form one of the cantilever arms, then the crack would have been visible and these parameters could have been observed as a function of temperature. Also, if the copper had been reacted to form oxide or other ceramic material, a measurement of a graphene-ceramic interface fracture resistance would be possible.
Another step forward was the realisation that the interface fracture energy could increase significantly with the ceramic processing conditions [28]. Using a nano-scratch adhesion measurement method, which is easy but unreliable, the as-transferred graphene adhered to SiO 2 with R= 3Jm -2 , a high value. But after rapid temperature rise and vacuum annealing, R rose to   A further area where graphene may give benefits is that it can be introduced as a thicker interphase layer, that might itself be toughened. This would limit the extent of crack growth along the laminae, which in the early materials, could reach almost to the ends of the sample.
The review [51] in 2013 showed that encouraging property improvements were being achieved with graphene in various ceramic matrices, especially electrical properties shown in Fig 6 [52]. The specific benefits were dependent on the graphene source and preparation, the mixing and compaction methods, then the sintering techniques. In this section we consider the bulk manufacturing processes. Then the next section deals with special processes for making functional ceramics, especially electrical components.
Bulk ceramic composite investigations have largely used graphene made from graphite, either by mechanical delamination, by the Hummer chemical splitting method to form graphene oxide, or by the electrolytic method. The largest demand is for polymer composites [53] but the same graphene preparations will also be applicable to ceramic products. High shear exfoliation using new designs of machine plus ultrasonic treatment has been of great interest [54]. Combined with the electrolytic method ( fig 7) using chemical additives to get between the graphene layers [55], followed by centrifuging to separate the nanoparticles, Hot pressing and isostatic hot pressing have been used to produce dense samples quickly [57].

Systematic Design of Nanoparticle Interfaces to Improve Properties
Mixing graphene nanoparticles with ceramic powders, then sintering the products, has shown significant but not stunning property improvements. The challenge is to design laminated structures which can give better results than random mixing, as pioneered in [5]. A recent embodiment of this idea is given by Belmonte et al [59]. They were striving to obtain    [62]. To deposit a high-k dielectric using ALD, it is necessary to introduce seed materials onto graphene due to the chemically inert surface of graphene or to generate seeding centers on graphene itself. These approaches cause heterogeneous dielectric stacks (or interfaces) and give rise to difficulty in controlling the film thickness, thereby constraining the scaling of gate dielectric thickness. A novel approach for the deposition of gate dielectrics therefore should be explored to achieve a single component gate dielectric that forms a homogeneous interface without the application of additive seed layers. One example would be the deposition of ultrathin (less than 10 nm) dielectrics by the initiated CVD method. The use of ultrathin dielectrics in graphene FETs would also be desirable for the development of flexible electronic devices. In summary, interface control is vital both for crack control and electronic applications. Several reviews have mechanical bias, but the electronic nature of the interface must be closely considered [64].

Conclusions and Forward View
Ceramics have been getting tougher over the past century as the mechanism of cracking processes around composite interfaces has begun to be But the most exciting challenge is to produce graphene interfaces in laminated ceramic composite structures, such that fracture energy can be increased by several orders of magnitude, producing much tougher and more multifunctional products than seen hitherto. The realistic prospect is that graphene reinforced ceramic composites could achieve toughness near R= 10,000 Jm-2 and K 1c = 20 MPam 1/2 at low graphene volume fractions around 1%, together with the possibility of novel electronic properties stemming from the precise interface control.