Scratching an Al/Si Interface: Molecular Dynamics Study of a Composite Material


We study scratching of a composite built from two widely different materials, a ductile and soft metal (Al) and a hard and brittle ceramic (Si). When scratching far away from the interface, the response of the pure elemental materials is monitored. A higher hardness and a lower friction coefficient are found for Si as compared to Al. The pile-up in Al is larger than in Si. When scratching along the interface, the composite responds approximately with the averaged behavior of the two pure materials. This applies to the forces as well as to the hardness and the friction coefficient. However, we observe a peculiar material flow, which can be described by a rotation around the scratch direction, inducing material mixing both in the groove bottom and in the pile-up; the harder Si expands on the groove bottom, while the softer Al expands in the pile-up region. When scratching across the interface, the material response switches on a length scale of the order of the contact radius. While the friction coefficient and the contact pressure switch from and to the values of the respective pure material, the behavior of the forces and areas is more complex. This is in particular due to the lateral pile-up that forms differently on the ductile metal and the amorphized Si parts.


Al/Si composites are relevant for today’s applications due to their light weight and are used for instance in the manufacturing of lightweight engine components [1,2,3]. The hard Si component serves to optimize the otherwise poor wear resistance of Al [1]. The originally used large (~ 100 μm) Si reinforcements created difficulties in machining of the Al/Si composite, such that later finer (2–5 μm) Si particles were used; the production of Al–Si composites by a spray-cast method allowed these composites to be more easily machined and to maintain an acceptable wear resistance [1]. Nowadays, the method of selective laser melting has been developed to produce Al–Si composites with extremely fine microstructure. These composites display both the least wear rate and the minimum wear loss compared with conventional castings [4].

Usually, the silicon particles stand out above the surface, and thus bear the load and provide scuffing and wear resistance during use, which discourages plastic deformation of the matrix [5,6,7,8]. For relatively short sliding time and low normal load, the wear rate is small. This stage is denoted as the ultra-mild wear stage [9]. At high normal load, the metal matrix is exposed to scuffing and deformation, which leads to high wear rates, denoted as mild wear [1]. It has been proved that the wear resistance of Al–Si composites both in the ultra-mild [2, 10] and in the mild wear regimes [1] is sensitive to the related microstructure including particle size and morphology [1], aspect ratio [10], and Si content [2].

Until now, experiments on the machining behavior of Al–Si composites only provide averaged results where the composite properties are summed over the individual behavior of Al and Si. In particular, no studies have been performed that identify the influence of the Al/Si interface on the mechanical behavior even though the interface can be expected to cause remarkable differences in pile-up, normal, and lateral load and friction coefficient as well as dislocation evolution and phase transformation. Therefore, up to now no clear answer has been obtained for the question how the Al/Si interface influences the machining behavior, for instance in scratching.

Atomistic simulations based on the method of molecular dynamics (MD) are increasingly being used to clarify the processes occurring on the nanoscale in machining of materials. Here the most intensely investigated process is nanoindentation [11], where a spherical indenter is pressed into a workpiece to test its response; data were obtained on the material hardness, but also on the processes of dislocation nucleation and growth in fcc [12,13,14], bcc [15, 16] and other [17, 18] metals and on phase transformations in Si [19, 20]. MD simulations of indentation in composite materials are more rare [21], and only indentation into multilayered films has been investigated more widely [22, 23]. Other machining processes, such as scratching and cutting, have been studied by MD only for homogeneous materials, such as metals [11, 16, 17, 24, 25] and Si [26,27,28]. We are not aware of any machining simulations in composite materials.

In the present study, we use MD simulations to investigate the fundamental mechanisms of scratching of Al–Si composites and clarify the influence of the interface on the response to machining.


Fig. 1

Setup of the simulation systems: a scratch along the interface; b scratch from Al to Si; c scratch from Si to Al. Units provided in nm. Green: Al; red: Si; white: initial position of indenter tip

We use MD simulations to study scratching of an Al/Si interface. Our simulation system consists of an Al/Si bicrystal, in which the {100} faces of a face-centered-cubic (fcc) Al crystal and of a cubic-diamond (cd) Si crystal form the interface. We also orient the top surfaces of the crystals in {100} direction.

We choose the size of the simulation crystals large enough that all dislocations generated are contained in the simulation volume. Since we perform three kinds of scratching processes—along the interface, across the interface from the Si to the Al part, and in opposite direction—we prepare three simulation systems; their sizes are sketched in Fig. 1. Note that the Al part needs to be larger than the Si part, since dislocations are created in fcc metals more abundantly than in Si. The thickness of our bicrystal—extension in z direction, see Fig. 1—amounts to 23.8 nm.

At the interface, the (100) face of the cd Si lattice has to be joined with the (100) face of the fcc Al lattice. Since Al and Si have different lattice constants—4.05 and 5.43 Å, respectively—this joining has to be done carefully. We follow the recipe of Noreyan et al. [29] and choose the lateral sizes of the interface (in y and z directions) such that they are commensurate for both the Al and Si lattice. In a previous publication [30], the quality of this interface has been characterized in detail; interface dislocations appear which periodically modulate the interface structure in \([01\bar{1}]\) direction.

The interactions between atoms in these three systems are described by the Al–Si potential developed by Saidi et al. [31]. It combines an embedded-atom-model (EAM) potential developed by Mendelev et al. [32] with a modification of the Stillinger–Weber potential [33] for Si. The ability of this potential to describe plasticity has already previously been tested, both for pure Al [34], pure Si [35,36,37,38], and for Al/Si composites [21, 21, 39].

Before starting the simulations, the systems are relaxed such that all components of the stress tensor reach values < 9 MPa and the temperature drops below 1 mK. This low temperature was used in order to ease the detection of dislocations in the scratched material. During scratching, we fix the bottom layers of all substrates with a width of 1 nm in order to prevent any rigid-body motion of the substrates. The next 1-nm layers at the bottom as well as the outermost 1-nm layers at all sides of the substrate are thermostatted to 0 K using velocity scaling. The boundaries in x direction are free, while periodic boundary conditions are applied in y direction.

The scratch tip is of spherical form with a radius of \(R=6\) nm. The interaction between the tip and an atom is purely repulsive and obeys the potential

$$\begin{aligned} V (r) = k (R-r)^3, \quad r \le R, \end{aligned}$$

where r is the distance between the atom and the center of the tip; for \(r>R\), the force vanishes. We use a force constant k with value of \(k=11.7\) eV/Å\(^3\) [21]. During indentation, the tip moves perpendicularly down into the surface to a final depth of \(d=4\) nm; during the scratch process, the tip moves parallel to the surface in a \(\langle 100 \rangle\) direction at a depth of \(d=4\) nm. The contact radius of the indenter hence amounts to \(a_{\rm{c}}= \sqrt{R^2-(R-d)^2}=56.6\) Å. The tip moves with a velocity of 20 m/s. It is known [34, 40, 41] that the effect of the tip velocity on the simulation results is small as long as it is below around 100 m/s; only at higher velocities, the substrate becomes heated by the tip-surface friction, and the results might be influenced.

For the determination of contact pressures, we need to calculate the contact areas of the tip with the substrate atoms; they are determined by summing the atomic areas of the so-called contact atoms. We identify the contact atoms by the procedure described in Ref. [25]: atoms are considered to be in contact with the indenter tip if their distance to the center of the tip is less than \(R+r_{\rm{c}}\). The material-dependent cut-off distances \(r_{\rm{c}}\) are calculated by comparison with the well-understood case of indentation, where the contact area can be safely calculated from a circular approximation [14]; we obtain \(r_{\rm{c}}=1.58\) (0.40) Å for Al (Si). The value for Si is so small since the high contact pressures push atoms close to the indenter surface. The atomic areas, \(A_{\rm{atom}}\), are determined from the atom number densities of the materials and amount to 7.80 (8.88) Å2 for Al (Si). In detail, the (projected) normal and tangential contact areas are given by

$$\begin{aligned} A_{\rm{norm}}= A_{\rm{atom}} \sum _i \cos \alpha _i, \quad A_{\rm{tang}}= A_{\rm{atom}} \sum _i \sin \alpha _i\,\cos\theta _i, \end{aligned}$$

where the summation is over all contact atoms i, \(\alpha _i\) denotes the angle (towards the surface normal) under which contact atom i is seen from the center of the tip, and \(\theta _i\) is the azimuthal angle under which atom i is seen with respect to the scratch direction.

The simulations are performed using the open-source code LAMMPS [42]. The data are analyzed with OVITO [43] and the Crystal Analysis Tool (CAT) [44,45,46].


Scratching Along the Al/Si Interface

Fig. 2

Groove and pile-up created during scratching along the interface. Snapshots have been taken after scratching a length of \(L=0\), 50, 90, 130, and 180 Å. Atoms are colored according to their height above the original surface; grayscale code is used for Si atoms and rainbow-scale for Al atoms

Table 1 Averaged scratching characteristics for scratching parallel and perpendicular to the interface

We first discuss the processes occurring when scratching along the interface. The response of Si and Al atoms during scratch is shown in the form of a top view in Fig. 2, which allows to discuss the formation of the scratch groove and the accompanying pile-up. Here it is seen that immediately after indent, \(L=0\), Si atoms in the indent pit have expanded across the interface to the Al side, while a few Al atoms in the pile-up forming around the indent pit moved onto the Al side. This behavior has been found previously and is caused by the higher compliance and smaller hardness of Al as compared to Si [21]. This behavior becomes even more pronounced during continued scratching. In particular, the upper part of the pile-up on the Si side is decorated by a top Al film. Since scratching is performed in a \(\langle 100 \rangle\) direction, and slip occurs both for the fcc Al and cd Si lattice structures in \(\langle 110 \rangle\) directions, the pile-up does not only form in front of the tip but also—and to a larger extent—sideways. Slip patterns in the Al surface become visible which move material sideways from the growing groove. The rim on the Si side consists of amorphized material [26], which expands sideways on the Si terrace. The Al material moves towards the Si side such that in particular the upper part of the pile-up on the Si side is decorated at the top by Al atoms. Thus, a sort of rotational motion is established where in the groove the hard material (Si) moves towards the soft (Al) side, while in the pile-up the soft material moves towards the hard side. The nature of this rotational motion—where atoms below the groove move towards the Al side, and atoms in the pile-up move towards the Si side—is further clarified in Fig. 3, which shows the atomic displacements of a vertical line of atoms during the scratch.

Fig. 3

Atom motion during scratching along the interface, shown in a cross-sectional view. The displacements of two selected groups of atoms on the Si and the Al side a are visualized when the scratch tip is at a distance of b 5 Å in front of and c 65 Å behind the displayed plane. The kink in the trajectories in c marks the positions the atoms had in b

During scratching, we monitor the values of the normal force, \(F_{\rm{norm}}\), the tangential force, \(F_{\rm{tang}}\), the friction coefficient, \(\mu\), the contact area, and the contact pressure as a function of scratching length. The results are displayed in Fig. 4; average values are assembled in Table 1. When calculating the average, the first 20 Å are not included, since these constitute an onset regime, in which the indentation stage continuously changes to the scratching stage. We note that we experimented with different values of the length of the onset regime; the values of the averages changed only negligibly. The normal and tangential forces are calculated as the sum of all forces exerted on the substrate atoms in normal (z) and tangential (y) direction, respectively, cf. Fig. 1a. Immediately after indent, the normal force decreases; this occurs while the rear part of the tip loses contact with the substrate. Concomitantly, the tangential force builds up. After scratching a length of around 20 Å, the forces have stabilized around the values displayed in Table 1; the remaining fluctuations are due to local emission of dislocations. The forces allow us to calculate the friction coefficient,

$$\begin{aligned} \mu = F_{\rm{tang}}/F_{\rm{norm}}, \end{aligned}$$

displayed in Fig. 4b; it stabilizes at a value of around 0.52.

The contact area of the tip with the substrate in normal and lateral direction plays an important role in quantifying the substrate response to scratching. It is calculated based on the cross-sectional areas of the substrate atoms contacting the tip [14, 25]. In this study, we calculate the contact areas of the Si part and of the Al part separately, and then sum over these two parts in Fig. 4c.

Similar to the force, the normal contact area drops after scratching started, as the rear part of the tip loses contact with the substrate [17, 25]; after around 20 Å, the total normal contact area assumes a stable value. Note that the Al and Si contributions to the area are not equal, even though the tip scratches symmetrically through the interface; the Si atoms have a larger contact to the tip than the Al atoms. This is due to the asymmetric reaction of the substrate to the scratch tip, which causes Si to expand into the Al part of the substrate, cf. Fig. 2.

The normal contact area is about twice as large as the tangential contact area, since the tip (radius of \(R=6\) nm) has been immersed only to \(d=4\) nm into the substrate.

During scratching, two contact pressures can be calculated from the normal and tangential material response, namely the normal and tangential pressures,

$$\begin{aligned} p_{\rm{norm}}= F_{\rm{norm}}/A_{\rm{norm}}, \quad p_{\rm{tang}}= F_{\rm{tang}}/A_{\rm{tang}}. \end{aligned}$$

The averages over these values are denoted as (normal and tangential) hardness, \(H_{\rm{norm}}\) and \(H_{\rm{tang}}\), during the scratch process.

Figure 4d displays the contact pressures. The normal contact pressure is remarkably stable throughout the scratch process. This is a sign that the indentation process has driven the material into the plastic regime such that forces and contact areas change in close proportion to each other. The lateral hardness again requires a scratch length of around 20 Å to reach saturation, since here the lateral forces need some scratch length in order to build up to their equilibrium values. As the averages in Table 1 show, the normal hardness (7.47 GPa) is slightly larger than the lateral hardness (6.66 GPa). This is in agreement with previous hardness studies in metals such as bcc Fe [17] and hcp Mg [25]; the resistance of a material to normal forces is larger than to lateral ploughing motion.

We can compare the quantities calculated here with those occurring when indenting an Al/Si interface. In Ref. [21], such an indentation study was performed with identical indentation conditions. At a depth of \(d=4\) nm, the normal force amounted to around 0.92 μN, larger than during scratch (0.60 μN); this is plausible since during scratch the entire circular of contact area of the tip is under load, while only (roughly) half of it is in contact with the substrate during scratch. However, the pressure during indentation, 7.5 GPa, agrees well with the normal hardness measured during scratch, 7.47 GPa. We conclude that our scratching results are in good agreement with previous indentation studies.

Fig. 4

Scratching characteristics as measured during scratch along the interface: a normal and tangential force; b friction coefficient; c contact area; d contact pressure. Black (gray) curves denote quantities measured normal (tangential) to the surface

Scratching Across the Al/Si Interface

Fig. 5

Groove and pile-up created during scratching across the interface from Al to Si. Snapshots have been taken after scratching a length of \(L=0\), 60, 120, 180, and 240 Å. Atoms are colored according to their height above the original surface; grayscale code is used for Si atoms and rainbow-scale for Al atoms

Fig. 6

Groove and pile-up created during scratching across the interface from Si to Al. Snapshots have been taken after scratching a length of \(L=0\), 60, 120, 180, and 240 Å. Atoms are colored according to their height above the original surface; grayscale code is used for Si atoms and rainbow-scale for Al atoms

The evolution of the pile-up when scratching from the Al side to the Si side is shown in Fig. 5. On the Al side, scratching develops a symmetrical frontal pile-up. Upon approaching the interface, the interface shows a slightly curved wavy structure. This is in part due to dislocation slip in Al, which pushes Al atoms on top of the Si part, see, e.g., the snapshot, Fig. 5, at 60 Å scratch length. A closer inspection of the subsurface structure of the Si block shows that it has been compressed elastically in front of the approaching scratch tip; due to high hardness of Si, no plasticity develops before the tip reaches the Si part itself, and the Si structure in front of the scratch tip is quite undisturbed. After passing the interface, Al atoms appear mixed into the surface of the groove. While the composition of the pile-up changes from Al to Si atoms, some Al atoms are present on the rim of the pile-up even after reaching the final scratch length.

When scratching in the opposite direction, Si → Al, the tip exerts a higher influence on the material in front, see Fig. 6. After traversing the interface, \(L=120\) Å, discrete slip steps in the Al material appear, which testify the evolution of an extended dislocation system in the Al part beneath the groove. The groove bottom is full of Si material, which has been transported a considerable way into the Al side. This transport is possible because Si is harder than Al and hence relaxes the pressure induced by the onward moving tip by expansion to the Al side. Note that this material covers in particular the left-hand side; this asymmetry in the evolution of the groove has been caused by a slight asymmetric fluctuation in the evolution of plasticity in the Si side. Since Si is considerably harder than Al, slight asymmetries in the Si response are magnified when the stresses reach the Al part. This asymmetry triggers the evolution in the Al part where a majority of slip and pile-up are created on the left-hand side.

Fig. 7

Cross-sectional side views of the damage created when scratching across the interface a from Al to Si and b from Si to Al, after a scratching length of 240 Å. Only a part of the entire simulation volume is displayed. Colors denote local lattice structure; in Si: gray (cd), dark (amorphous), light gray (atoms neighboring dislocations and stacking faults as well as interfaces); in Al: green (fcc), blue (disordered and interface atoms), red (hcp)

We provide a side view of the scratched composite in Fig. 7, which allows us to discuss the depth at which damage has been created. The local lattice structure has been identified using the method of common neighbor analysis [47, 48] and its recent extension to cd structures [49] as implemented in OVITO [43]. The amorphization of the groove bottom and of the pile-up in Si are clearly seen; in addition, several dislocations have been created in front of and beneath the groove. Also the surface and interface of the Al part appear disordered in Fig. 7. Stacking faults created in Al can be identified by their local hcp lattice arrangement. These stacking faults are created by dislocation activity, which will be analyzed further in Sect. 3.3. We conclude that the damage extends considerably deeper in the ductile metal than in the brittle Si part.

Quantitative data on the results of scratching across the interface are displayed in Figs. 8 and 9; again the averaged data are assembled in Table 1. Since we studied both scratching from the hard material (Si) towards the soft material (Al) and vice versa, the number of data has doubled in comparison to the scratch along the interface. Note that scratching always starts 120 Å in front of the interface and ends 120 Å after it; if we plot the data as a function of scratch length L, the interface is hence always situated at \(L=120\) Å, in the middle of the graphs. The averages in Table 1 are again taken without including the onset regime of the first 20 Å, where the indentation stage passes towards the true scratching stage. In addition, a region extending 20 Å in front of and behind the interface is excluded, such that the data pertain to scratching in the undisturbed Si and Al parts only. Again, we experimented with different values of the length of the excluded region around the interface; the values of the averages changed only negligibly.

The normal force, Fig. 8a, shows a simple behavior when scratching from the Al side to the Si side. After a short onset regime of 20 Å length, the force stabilizes at a low value of 0.45 μN. While traversing the interface, the force rises steadily and stabilizes at the high value of 0.82 μN; this difference is in agreement with the higher hardness of Si as compared to Al. For the opposite scratch direction, the behavior is different. After indentation into the Si part and starting scratching, the normal force steadily decreases and does not saturate. Here, the presence of Al softens the material response of Si even at such large distances as 120 Å from the interface, such that the normal force of pure Si—which must be in the region of 0.82 μN, as determined in the Al → Si scratch—does not stabilize. After traversing the interface, the normal force stabilizes at 0.42 μN, a similar value as in the Al → Si simulation (0.45 μN). The size of switching region is of the order of the contact radius, \(a_{\rm{c}}=57\) Å. The switching region appears to be somewhat smaller, since the changes do not start suddenly when the front part of the tip reaches the other material, but more gradually.

The evolution of the lateral forces with scratch length is displayed in Fig. 8b. These curves show much less dependence on the material that is scratched and on the scratch history than the normal forces. Apart from the onset regime, the values are well (within roughly 25%) described by the reference value given by scratching along the interface. In detail, Table 1 shows that the average lateral force on the Al side (0.28 μN) is smaller than on the Si side (0.33–0.36 μN). However, the difference of the materials is smaller than for the normal forces; this is due to the relatively small friction coefficient of Si to be discussed below. The data on the Si side show again a more pronounced dependence on the scratch path; after traversing Al, the lateral force is smaller in Si than when starting in Si. This again demonstrates that the admixture of Al to the pile-up has a softening effect on Si.

The friction coefficient, Fig. 8c, shows a comparatively simple behavior in that—apart from an onset regime in Al—its value mainly depends on the material beneath the tip and appears to be quite independent of the path taken. The friction coefficient on the metal Al (\(\mu =0.65\)–0.68) is larger than on the hard Si (\(\mu =0.43\)–0.44). The values for Al coincide well with those measured previously [50] in pure Al for the same (001)[010] scratch system, and with similar scratch conditions (indenter radius \(R=5\) nm, scratch depth \(d=3\) nm); there a friction coefficient of \(\mu =0.61\) was obtained. We are not aware of comparable data for scratching single-crystalline Si. The lower values obtained by us are, however, in line with the general finding that hard materials exhibit smaller friction coefficients than soft materials [51].

Note that on the location of the interface, \(L=120\) Å, all values of the forces and of the friction coefficient roughly coincide with those for scratching along the interface. This demonstrates that scratching along the interface provides a response that is roughly the average of the pure elements.

Fig. 8

Scratching characteristics as measured during scratch across the interface: a normal force; b tangential force; c friction coefficient. Green: scratch from Al to Si; red: scratch from Si to Al. The horizontal line provides the average value for scratching along the interface as a reference

Fig. 9

Scratching characteristics as measured during scratch across the interface: a normal contact area; b tangential contact area; c contact pressure. Green: scratch from Al to Si; red: scratch from Si to Al. The lighter (darker) curves in c denote tangential (normal) pressures. The horizontal line provides the average value for scratching along the interface as a reference

The normal contact area, Fig. 9a, is—after the onset regime—well described by the average area determined for scratching along the interface. An interesting issue is provided by the fact that, when scratching from Si to Al, the contact to atoms of the original species (Si) stays for a longer time and the contact to the new species (Al) is delayed as compared to the opposite scratch path. This behavior is again due to higher hardness of Si which let this material expand into the Al part.

The lateral contact area, Fig. 9b, shows stronger deviations from the behavior for scratch along the interface that we take as a reference. Note in Table 1 that the lateral contact area of Al is systematically larger than that of Si; this is a feature of the softer and more ductile behavior of Al which produces higher pile-ups. When scratching from the Al side towards Si, the contact area in Al is increased with respect to this reference; this is caused by the larger ductility of Al. When entering into the Si side, the lateral contact area, however, well follows this reference. However, when starting from the Si side, the lateral contact area steadily increases until passing the interface, and only then saturates. The increase in lateral contact is then mainly due to Al atoms, which create an increasingly high pile-up, cf. also the snapshots provided in Fig. 6. Note that there the color bars were modified in order to represent the highest pile-ups that reach 80 Å; these constitute the highest pile-ups generated in this study.

Figure 9c displays the normal and lateral contact pressures. When scratching from the Al to the Si side, a simple behavior is seen in that the contact pressures change at the position of the interface and have stable values before and behind the interface. In detail, the normal and tangential values on the Al side roughly coincide, while the tangential pressure on the Si side is smaller than its normal counterpart. As Table 1 demonstrates, this is a systematic feature, which we attribute to the fact that Si yields not so much by dislocation plasticity, cf. Sect. 3.3, as by amorphization. Amorphous Si develops more easily near the surface, in the pile-up, than in the substrate interior [28]. In consequence, the resistance of Si to lateral motion—where material may evade the pressure towards the surface—is smaller than to normal forces, where such a remedy is not available. When scratching in the opposite direction, a roughly symmetrical behavior results; in particular, the values of the hardness depend only negligibly on the scratch direction, Table 1. Note, however, that again the Si \(\rightarrow\) Al scratch direction leads to a longer survival of the high Si hardness values; this behavior is caused by the expansion of Si material into the Al part during scratch.

The force plots, Fig. 8a and b, and the pressures, Fig. 9c, show an interesting asymmetry in that the switch from the high values on the Si side to the low values on the Al side is not centered of the interface itself (scratching length 120 Å), but always in front of it. So for instance, for the normal force, Fig. 8a, when starting from the Al side, the normal force increases upon approaching the interface; however, the average value is already encountered at around 100 Å. This occurs analogously when approaching the interface from the Si side. This feature is due to the fact that we measure the position of the indenter with respect to its center; however, during scratching, only the front hemisphere of the indenter is in contact with the substrate. Since the approach of the interface is felt ‘earlier’ by the front part of the tip than the central part, the tip response starts already before the tip center is positioned on the interface.


Fig. 10

Length of dislocation lines created when scratching a along the interface, b across the interface from the Al to the Si side, and c across the interface from the Si to the Al side. Note the two different ordinate scales, one referring to dislocations in Al (circles), and one to dislocations in Si (squares)

Indentation and scratching generate plasticity in the material, which is based on dislocations in the case of metals [34, 41, 50]. In the case of Si, amorphization of the lattice structure around the indent tip occurs [28], but also dislocations are generated, albeit in smaller numbers [20]. Amorphization of the groove and generation of dislocations in its environment have been demonstrated in Fig. 7. We quantify the amount of dislocations created by identifying the dislocation lines in our simulation volume at the end of the scratch; the length of each dislocation line is determined using OVITO [43] and summed up to obtain the total length of dislocation lines in the Al and the Si side. This quantity is displayed in Fig. 10; note that different ordinate scales have been used for Al and Si, which demonstrate that an order of magnitude more dislocations are created in Al as compared to Si.

When scratching along the interface, Fig. 10a, a very different behavior of Al and Si shows up. On the Al side, the length of dislocations steadily increases; this behavior is typical for fcc materials, which show ample dislocation activity under scratch [50]. In Si, however, the length of dislocations stays roughly at the same value as that reached after indentation. We presume that the reason for this behavior is the Al/Si interface; Si can relax the high pressure induced by the scratch tip by expansion into the nearby Al material [21]; due to the pressure relaxation, the creation of further dislocations is inhibited.

Scratching across the interface, Fig. 10b and c, exemplifies again the different response of Al and Si. When scratching from the Si to the Al side, the length of Al dislocations increases almost linearly with scratch length, as soon as the interface has been approached by the scratch tip—in fact already a distance of 60 Å in front of it, corresponding to the contact radius. When scratching from the Al side towards the interface, Si dislocations are only generated when the tip approached to less than 20 Å to the interface; this demonstrates the huge pressures necessary for dislocation generation in this material. The length of Si dislocations then does not increase smoothly but rather jumps up suddenly. The values reached then (around 600–800 Å) are of similar size as the dislocations left behind when the scratch started on the Si side; it is considerably larger than the value of 200–330 Å for scratching along the interface. This demonstrates again that the vicinity of the soft Al prevents dislocation build-up in Si for the scratch geometry along the interface.


We performed a molecular dynamics simulation of scratching of a composite material consisting of two widely different materials, a ductile and soft metal (Al) and a hard and brittle ceramic (Si). By scratching both along and across the interface, we could study the response of the composite to this machining process. We obtained the following major findings.

  1. 1.

    When scratching far away from the interface, the response of the pure elemental materials is established. A higher hardness and a lower friction coefficient are found for Si as compared to Al. The pile-up in Al is larger than in Si, which can be attributed to the softer and more ductile character of the metal.

  2. 2.

    While in the pure Al part, the tangential hardness is almost as large as the normal hardness, in Si the tangential hardness is 20 % smaller than the normal hardness. We attribute this feature to the amorphization of the material in front of the tip, which makes the resistance to lateral motion smaller than to normal forces.

  3. 3.

    When scratching along the interface, the composite responds approximately with the averaged behavior of the two pure materials. This applies to the forces as well as to the hardnesses and the friction coefficient.

  4. 4.

    In the groove the Si material expands towards the Al side. On the other hand, in the pile-up Al atoms are transported onto the Si side. These atom movements correspond to a rotational atom flow around the scratch direction, inducing material mixing both in the groove bottom and in the pile-up.

  5. 5.

    When scratching across the interface, the material response switches on a length scale of the order of (but smaller than) the contact radius.

  6. 6.

    While the friction coefficient and the contact pressures switch from and to the values of the respective pure material, the behavior of the forces and areas is more complex. This is in particular due to the lateral pile-up that forms differently on the ductile metal and the amorphized Si parts.

  7. 7.

    While in all cases the length of dislocations generated in Si is more than an order of magnitude smaller than in Al, Si dislocations are even more suppressed when scratching along the interface. In this case the pressure relaxation by Si expansion to the Al side reduces the nucleation and growth of Si dislocations.


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Simulations were performed at the High Performance Cluster Elwetritsch (RHRK, TU Kaiserslautern, Germany). We acknowledge the financial support of the Deutsche Forschungsgemeinschaft via the IRTG 2057 and the SFB 926.

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Correspondence to Herbert M. Urbassek.

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Zhang, Z., Alabd Alhafez, I. & Urbassek, H.M. Scratching an Al/Si Interface: Molecular Dynamics Study of a Composite Material. Tribol Lett 66, 86 (2018).

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  • Metal matrix composites
  • Atomistic simulation
  • Scratching
  • Al/Si interface
  • Hardness