B D Beake

V M Vishnyakov

J S Colligon

Nano-impact testing of TiFeN and TiFeMoN films for dynamic toughness evaluation

Keywords: nano-impact, nanoindentation, wear resistant coatings, thin films

TiFeN and TiFeMoN films were deposited on silicon wafers by ion beam assisted deposition. Their mechanical properties were measured by nanoindentation (quasistatic) and nano-impact (dynamic) techniques. Nano-impact testing enabled assessment of their toughness and resistance to fatigue fracture under repetitive loading. At low impact forces films with a higher resistance to plastic deformation (H 3 /E 2 ) were much more resistant to the formation of cracks throughout the test. At higher impact forces these films initially show impact resistance but with continued impacts they are unable to protect the Si substrate, performing as poorly as films with lower H 3 /E 2 and suffer delamination from the Si substrate over a large area.

Introduction

Industrial applications such as high speed cutting and forming processes have increasing severe mechanical contact conditions that drive the development of materials with a combination of hardness and toughness [START_REF]Physical and Mechanical Properties to characterize Tribological Compatibility of Heavy Loaded Tribo-systems[END_REF][START_REF] Beake | Using nanomechanics to optimise coatings for cutting tools[END_REF]. The optimum mechanical properties required for film toughness remain an open question. From a practical viewpoint, "super-tough" can be preferable to "super-hard" in highly loaded mechanical contact [START_REF]Physical and Mechanical Properties to characterize Tribological Compatibility of Heavy Loaded Tribo-systems[END_REF][START_REF] Beake | Using nanomechanics to optimise coatings for cutting tools[END_REF]. Nanoindentation has been used to provide an estimate of toughness [START_REF] Lawn | [END_REF][4][5][6][7]. The quantitative treatment developed by Lawn and co-workers for bulk materials [START_REF] Lawn | [END_REF][4] is not directly applicable to thin films as the presence of the interface influences the development of the radial crack system. For highly brittle thin films when discrete through-thickness fracture is observed in the load displacement curve (e.g. 1µm CN x on Si) modelling may provide quantitative estimates of toughness [5][6][7]. However, such fracture is not observed in the load-displacement curve of tougher thin films. Chen and Bull have noted that for some coatings, even when fracture does occur, e.g. picture frame cracking of ZnO, there are no excursions on the load-displacement curves [8]. There are opposing views about how the indentation curve shape should look to correspond to a tougher material, when discontinuities (abrupt depth jumps) due to cracking are not present, and parameters such as the ratio of hardness to modulus (H/E, a measure of the partitioning between energy dissipation by elastic and plastic deformation in the indentation contact) [9][10], plasticity index (plastic work done/total work done) [11][12] and the resistance to plastic deformation (H 3 /E 2 ) [13][14][START_REF] Musil | Nanostructured coatings[END_REF][START_REF] Musil | [END_REF][17][18][19][20] have been investigated as these can be easily obtained in a nanoindentation test. The shape of the indentation curve itself can be considered as a "fingerprint" for a material as it is directly related to the H/E r ratio in the contact, where E r is the reduced indentation modulus as defined in Equation 1 (see section 2). An increase in H/E results in a lower plasticity index and a more elastic contact. As an illustration nanoindentation curves for films with differing plasticity are shown in Figure 1. The plasticity index for the indentation on the softer sample (squares) is 0.48 whilst for the harder sample (circles) it is 0.31.

Toughness has been associated with minimisation of plasticity in indentation, particularly in the development of super-hard nanocomposite materials [21][22]. For hard elastic films almost complete recovery occurs in indentation and the plasticity is minimised. Musil and Jirout [START_REF] Musil | [END_REF] have shown that, for indentation load up to 0.5 and 1N with a Vickers indenter, resistance to cracking is correlated with the film resistance to plastic deformation, H 3 /E 2 . In Figure 1 the measured H 3 /E r 2 values are 0.18 and 0.64, and, based on the Musil and Jirout analysis, the film with less plasticity would have higher resistance to cracking. The open question is whether this is beneficial in a realistic practical contact situation. When coatings are developed for highly loaded contact applications [START_REF]Physical and Mechanical Properties to characterize Tribological Compatibility of Heavy Loaded Tribo-systems[END_REF][START_REF] Beake | Using nanomechanics to optimise coatings for cutting tools[END_REF][START_REF] Lawn | [END_REF][11][12][23][24][25][26][27][28] higher plasticity in indentation has been correlated with higher toughness and durability in application conditions.

Toughness is not the same as plasticity, but under highly loaded conditions cracks can propagate to failure where dissipative mechanisms such as plasticity do not exist so the two properties can be very well correlated.

In view of this apparent contradiction regarding the optimum mechanical properties for toughness, in this study a repetitive nanomechanical test has been employed to take a more direct approach to the assessment to film toughness and damage tolerance under dynamic loading. Dynamic loading represents a typical mechanical contact situation where high dynamic toughness is required to avoid coating failure problems.

The nano-impact (impulse indentation) technique in the NanoTest system can simulate the mechanical conditions occurring in contact and predict the wear performance of coated systems subjected to repetitive contact [START_REF] Beake | Using nanomechanics to optimise coatings for cutting tools[END_REF][29][30][31][32]. In this test repetitive high strain rate impact indentation occurs at the same location on the sample surface and the evolution of film damage can be monitored in situ. The test does not provide a direct value of the static fracture toughness K 1c but rather an assessment of the sample's resistance to fatigue fracture, or effective dynamic toughness, under dynamic and repetitive loading conditions that can be more representative of actual contact in applications. For example, the nano-impact test can successfully simulate the contact in extreme high speed cutting applications such as machining hard-to-cut materials with 3 µm AlCrN or TiAlN coatings deposited on cemented carbide tool inserts [START_REF] Beake | Using nanomechanics to optimise coatings for cutting tools[END_REF]12]. At the start of the impact test the penetration depth is lower for the TiAlN (higher H, H/E and H 3 /E 2 ) but as the test progresses it is more susceptible to fracture and this results in greater wear. This exactly mirrors the observed behaviour in high-speed cutting tests on hardened steel [START_REF] Beake | Using nanomechanics to optimise coatings for cutting tools[END_REF]12].

In this current study we investigate the behaviour of a range of hard brittle nitride films deposited on a hard and brittle but low modulus substrate (Si wafer) under continued impact at different forces. The primary aim was to determine how the film mechanical properties influence impact resistance: (1) does increasing film H 3 /E 2 increase toughness in dynamic contact? (2) are differences observed between film resistance to single impacts and prolonged repetitive impacting? (3) can the results estimate the optimum film mechanical properties necessary for improved toughness and (4) can this data help to reconcile the opposing standpoints on whether plasticity in indentation is desirable for toughness or not?

Experimental

The films were prepared on Si (100) wafers cleaned using the ion-assisting beam with Ar + ions. A dual ion beam sputtering system described in detail previously was used [38][39]. Although deposition rates are slower than in commercial PVD and CVD coating processes, the dual ion beam approach can provide the flexibility and control to deposit films with a wide range of mechanical properties [39]. Sputtering of the target was by a 1.25 keV Ar + ion beam and a 280-600 eV N 2 + ion beam was used for ion-assistance of the deposited film. The sputter target comprised either (i) Ti and Fe or (ii) Ti and Fe and Mo, with the area occupied by the Fe (or Fe and Mo)

determining the relative Ti-Fe(-Mo) ratio in the coating. An initial thin Ti layer was deposited first on the Si to improve adhesion and the hard coating then deposited onto this bonding layer. Nitrogen ions from the second source bombarded the growing coating during ion-assisted nitride sputtering which occurred in a nitrogen partial pressure of 1-2 x 10 -2 Pa (see Table 1 for details). The substrate temperature was maintained constant during all the deposition steps and the film thickness was determined using an SEM. Film microstructure has been assessed by X-ray diffraction (XRD) using a Philips X'pert 20 diffractometer with Cu anode in theta-2theta configuration.

Nano-impact testing was performed using a Micro Materials NanoTest NTX fitted with a sharp cube corner diamond indenter as an impact probe. The technique has been described in detail elsewhere [START_REF] Beake | Using nanomechanics to optimise coatings for cutting tools[END_REF][31][32][33][34][35][36][37]. In addition to the repetitive impact experiments described in this current work the technique can also be used for single impacts, obtaining properties such as dynamic hardness and coefficient of restitution [33][34][35][36]. The position of the impacting probe is recorded throughout the test in the instrument software. At the beginning of the test the surface is contacted by the impact probe under a minimum contact load (0.03 mN) to determine the depth zero.

The static coil force is then applied (5-50 mN) producing elastoplastic deformation by indentation. By solenoid activation the impact probe is rapidly removed to 12 µm from the surface and then accelerated over this distance to impact the surface causing additional deformation due to the higher inertial force. Subsequent repetitive impacts every 4 seconds result in additional deformation depending on the mechanical properties of the films. The probe was accelerated at applied forces of 5-50 mN for the test duration of 300 s. The evolution of impact damage was followed in situ by probe depth vs. time plots. Abrupt increases in probe depth during the tests correspond to through-thickness cracking. Post-test microscopic analysis of impact craters confirmed through-thickness cracking and resultant film delamination. Repeat tests were performed for each sample at each load to assess any variability. The standard deviations of the initial (pre-impact) depths are within the size of the symbols in figure 4a. At 25 and 50 mN, where fracture is appreciable, the standard deviation in the final impact depth (figure 4 c) is typically in the range 10-20 %.

Nanoindentation was also performed using the Micro Materials NanoTest NTX.

Multi-cycle load-controlled "load-partial-unload" experiments were performed up to 20 mN maximum load to assess the variation in mechanical properties with penetration depth. The unloading curves were analyzed using standard methods with the area function for the Berkovich diamond indenter determined by indentations into fused silica. Film-only hardness (H film ) was determined from the plateau (loadinvariant) region and film-only reduced elastic modulus (E r ) film by extrapolation to zero depth (extrapolation procedure according to ISO 14577-4) [START_REF]Metallic Materials -Instrumented Indentation Test for Hardness and Materials Parameters, Part 4 -Test Method for metallic and non-metallic coatings[END_REF]. E r is the reduced modulus defined by

1/ E r = (1-ν s 2 )/ E s + (1-ν i 2 )/ E i [Eqn. 1]
where ν s = Poisson's ratio for the sample, ν i = Poisson's ratio for the diamond indenter (0.07), E s = Young's modulus for the sample, E i = Young's modulus for the indenter (1141 GPa). Throughout this paper for simplicity we have reported properties

(such as H/E r and H 3 /E r 2

) based on E r rather than E since E r is directly measured in the nanoindentation test and does not require the film Poisson's ratio to be known or estimated.

Results

Nano-impact testing

The sample response to repetitive impact varied with film composition and impact force. As an example, the transition in failure behaviour on the TiFe 0.65 N 1.45 film surface after impact testing for 300s at 5, 10 and 25 mN is shown in Figure 2. At 5 mN the film is resistant to cracking but localised fracture occurs at 10 mN and at 25 mN both fracture and large-area film delamination are observed.

More information can be obtained from the displacement vs. time responses during the test. Typical nano-impact behaviour at each force is shown in Figure 3. At 5 mN, the TiFe 0.69 N 1.44 and TiFe 0.65 N 1.45 films are resistant to cracking throughout the test and the final depth is around 500 nm (figure 3 a). The TiFe 0.63 N 1.5 film shows the lowest impact resistance, and delamination occurs on the second impact. The TiFe 0.7 N 1.9 and TiFe 0.4 Mo 0.4 N 2 films show an impact behaviour between these extremes where an initial gradual increase in depth proceeds until the onset of a more marked deformation mode after 20 s for TiFe 0.4 Mo 0.4 N 2 and after 35 s for TiFe 0.7 N 1.9 .

At 10 mN the TiFe 0.69 N 1.44 and TiFe 0.65 N 1.45 films perform well with TiFe 0.65 N 1.45 being crack-resistant and TiFe 0.69 N 1.44 showing only a small increase in depth after 80 s (figure 3 b). For the other three films, cracking occurs almost immediately.

At 25 mN the TiFe 0.69 N 1.44 and TiFe 0.65 N 1.45 films are more resistant to cracking after initial impact but, upon continued impact, they fracture dramatically (figure 3 c).

More extreme substrate deformation occurs at 50 mN (figure 3 d), with all films fracturing rapidly and unable to protect the substrate from further damage. Final probe depths were in the range 5-10 µm.

Key parameters in the impact test are summarised in figure 4 a-d. These are

(1) probe depth prior to initial impact

(2) probe depth after first impact

(3) final impact probe depth (4) increase in probe depth due to continued impact (i.e. the difference in depth between 3 and 2 above) 1). The harder films showed enhanced stiffness but the value of (H/E r, ) film covered a relatively small range (0.083-0.093).

Nanoindentation

Film thickness varied between 0.7 and 1.7 µm. More marked differences were found in the parameter (H 3 /E r 2 ) film , a measure of the film-only resistance to plastic deformation which varied between 0.11 for the softest film and 0.21 for the hardest.

Figure 5 shows how the resistance to plastic deformation in the contact varies with increasing penetration into the films. Although the value of H 3 /E r 2 in the contact varies with penetration depth (primarily due to the increased contribution from the much lower modulus Si substrate), the data clearly show that TiFe 0.69 N 1.44 and TiFe 0.65 N 1.45 films are significantly more resistant to plastic deformation.

Microstructural analysis

Figure 6 shows the XRD spectra of the films. In conjunction with ion assistance, the high N 2 partial pressure during deposition of TiFe 0.69 N 1.44 and TiFe 0.63 N 1.5 resulted in fully amorphous films whilst the TiFe 0.69 N 1.44 and TiFe 0.7 N 1.9 deposited under lower N 2 partial pressure shows peak in their XRD spectra. The TiFe 0.4 Mo 0.4 N 2 film was amorphous.

Discussion

The films are compositionally different from Ti(Fe)N x films deposited by Musil and co-workers who reported that, when the Fe content was sufficiently low (≤15 at.%), single-phase nanocomposite films could be obtained with hardness up to 45 GPa [14].

Our films contain a much higher % Fe and are softer due to FeN formation. The XRD analysis has shown that the TiFeN film structure is controlled by both the nitrogen partial pressure and the presence of ion assistance. At high nitrogen partial pressure ion assistance produces fully amorphous TiFeN films as the temperature of the substrate is not high enough to recrystallise them. The TiFe 0.4 Mo 0.4 N 2 film is also amorphous, presumably as with more elements it is unlikely to be crystalline.

At ~20 GPa the TiFeN and TiFeMoN films have similar hardness to thin films (such as a-C, nc-TiC/a-C, TiC/a-C, TiC/a-C:H, nc-TiC/a-C(Al), YSZ/Au) reported to have excellent toughness [17][18][23][24][25]28] and high plasticity in indentation. The resistance to plastic deformation H 3 /E 2 and the partitioning between energy dissipation by elastic and plastic deformation H/E are well correlated for these thin films. Musil and coworkers have shown that nanocomposite films with higher hardness also tend to have higher H 3 /E 2 [13][14][START_REF] Musil | Nanostructured coatings[END_REF][START_REF] Musil | [END_REF] and this is consistent with the results in Table 1. In what follows the discussion has been based around H 3 /E r 2 for convenience but Table 1 shows that similar arguments also apply for H/E r .

Nanoindentation can be used to assess resistance to crack initiation and overload failure but it is not possible to monitor crack propagation under repetitive, oscillating loading conditions in the conventional quasi-static test. The dynamic nano-impact test is a complementary technique that can determine the resistance of the film-substrate system to fatigue fracture, i.e. the formation and propagation of cracks under repetitive loading. Tougher films with plastic deformation and micro-cracking mechanisms appear beneficial in minimising large-area delamination under these conditions. The abrupt increases in impact probe depth (such as those observed in test. This is sufficient to avoid cracking at 5 mN but lead to the propagation of small cracks causing total failure at 10 mN.

Examination of the data presented by Pei and co-workers shows that the onset critical load for indentation-induced radial cracking with a cube corner indenter correlates well with H 3 /E 2 for a range of 1.5 µm nc-TiC/a-C:H films [17]. Musil and co-workers have reported that H 3 /E r 2 and film thickness both influence cracking resistance in microindentation with a Vickers indenter, with thicker and higher H 3 /E r 2 films being more resistant [START_REF] Musil | [END_REF]. The low thickness of TiFe 0.63 N 1.5 and TiFe 0.7 N 1.9 films may be a factor in their low crack resistance at 5 and 10 mN. The onset of circular cracking has been correlated with the ratio of indentation depth to film thickness by Musil and coworkers in microindentation [START_REF] Musil | [END_REF]. In the current study the sputtering effect at high ion assistance results in more intense ion bombardment and smaller film thickness. At 700 nm, TiFe 0.63 N 1.5 is the thinnest of all the films and fractures immediately at 5 mN.

But film thickness is not the sole factor since one of the thickest films, TiFe 0.4 Mo 0.4 N 2 , also fractures at low load. Thickness has been shown previously to affect the nano-impact response of 2.3-4.5 µm a-C:H films on at low load (5 mN, cube corner indenter). Thicker a-C:H films showed more damage tolerance, providing load support for the substrate even after some cohesive fracture, although there was little effect on the number of impacts required for cohesive fracture. In general, film thickness has been found to influence nano-impact behaviour to a much smaller extent than nano-scratch testing. For example, critical loads for film failure in nanoscratching ta-C films on Si and a-C:H-based films on glass scale with film thickness [START_REF] Beake | [END_REF][42] but the number of impacts required for failure can show an opposite trend. In nano-impact tests at low load (5 mN, cube corner indenter) a 140 nm a-C:H with H 3 /E r 2 = 0.21 required almost twice as many impacts to fail as a 600 nm Si-doped a-C:H with H 3 /E r 2 = 0.16.

The impact results at 25 and 50 mN applied load (fig. 3 c-d) do not show the same correlation with the mechanical properties of the films observed at lower load. As the impact force increases the enhanced load support provided by the higher H 3 /E r 2 films is no longer sufficient to protect the substrate from deformation and causes film fracture and increasing impact-induced damage throughout the test. A sharp decrease in film durability with increasing applied load in the nano-impact test has been reported previously for 2.3 µm a-C:H/Cr on steel with H 3 /E r 2 = 0.21 [30]. At low load this coating was relatively durable. It showed only small cohesive fractures after long impact times but, above a threshold load of 10 mN, through-thickness cracking led to complete removal of the coating soon after the start of the test. Bousser and coworkers have shown that for 8-13 µm CrN and CrSiN films on stainless steel the ratio of indentation depth to film thickness (d/h) at which circular cracking in Vickers indentation occurs was correlated with plasticity and inversely correlated with

H 3 /E r 2 ,
suggesting that, under these highly loaded conditions, the inability of the coating to accommodate the substrate deformation drives the cracking process [19].

From individual impact depth vs. time graphs (as in figure 3) it is possible to separate film behaviour into three distinct regimes as the test progresses: (i) load support (preimpact) (ii) initial impact and (iii) subsequent impacts. This procedure enables the role of film properties on load-support, initial resistance to crack nucleation and subsequent crack propagation to be investigated in more detail. In figure 4 the data for two films with very high H 3 /E r 2 , (dotted lines) have been contrasted with the data for films with lower H 3 /E r 2 , (full lines). The initial load-support prior to impact is greater for these two films throughout the load range (fig 4(a)). The probe depth after a single impact shows a similar trend (fig 4(b)). It is an appreciable fraction of the film thickness even at low load. Figure 4 (c) shows that the increase in depth due to repetitive impact (i.e. final depth -depth after initial impact) varies with load. At low load this increase is smaller for the two TiFeN films with higher H 3 /E r 2 , since these do not fracture on repetitive impact. However, comparison of figures 4 (b) and (c) shows that, although resistance to single impact is improved by increasing H 3 /E r 2 , durability at high dynamic contact load is not. At lower impact load, the ratio of impact depth to film thickness is low so higher H 3 /E r 2 , in the film is beneficial by preventing cracks from forming in the first place. Increasing H 3 /E r 2 , has been associated with high resistance to cracking and hence high fracture toughness [13][14][START_REF] Musil | Nanostructured coatings[END_REF][START_REF] Musil | [END_REF][17][18][19][20]. At higher impact load, effective durability and damage tolerance requires resistance to both formation and propagation of cracks under dynamic loading. For Vickers indentation, it has been reported that cracks often form when d/h > 0.5 [19], so it is reasonable to assume cracks form on first impact at higher load, and the subsequent film behaviour in figures 3c and 3d is the response of increasingly cracked films to continued impacts.

The two TiFeN films with higher H 3 /E r 2 , have better initial impact resistance but as the test continues fracture occurs progressively and they are unable to protect the substrate from deformation. The brittle films cannot accommodate this deformation and tensile stresses arise causing film cracking leading to the dramatic delamination failure (as observed in figure 2). Dissipative mechanisms such as plastic deformation and micro-cracking may be beneficial in relieving accumulated strain for the films with lower H 3 /E r 2 . Large area delamination occurs in nano-scratch testing due to tensile stress accumulation on high H 3 /E 2 TiFeN and nc-TiN/a-Si 3 N 4 nanocomposites deposited on silicon [43] but is absent on films with lower hardness and H 3 /E 2 . Zhang and co-workers have shown that the durability nc-TiN/a-Si 3 N 4 nanocomposites in scratch tests was enhanced at hardness ~20 GPa, gradually decreasing as hardness approached 40 GPa [26].

The apparent disparity in the literature reports relating the shape of the indentation load-displacement curve to film toughness may well relate to the severity of the contact conditions under which toughness is assessed. Plasticity can provide a reasonable first estimate of toughness in the engineering sense as plasticity is related to fracture resistance. Fox-Rabinovich and co-workers have shown that the optimum mechanical properties for fracture resistance and tool life in cutting tool coatings depends strongly on the type of cutting test [11]. At low load where fracture is less likely, improving film wear resistance by minimising plasticity is more important.

Substrate deformation can then be prevented by the coating having a high H 3 /E 2 value so that cracking is minimised. Although the TiFeN films possess no intrinsic resistance to high load impact when deposited on the brittle and low modulus Si substrate, they may yet have the potential to be used on higher modulus substrates, with bias-grading for stress-relief. Ma and co-workers have shown that wear resistance of nc-TiN/a-Si 3 N 4 in pin-on-disk tests is improved by high hardness [22].

Film deposition with graded interface layers between the substrate and the composite coatings has also been shown to enhance the scratch test critical load and relieve interface stresses [24]. Guruvenket and co-workers have recently reported that surface nitriding of low alloy steel substrates prior to deposition of nc-TiN/a-SiN x and nc-TiCN/a-SiCN films produced a more gradual change in hardness from substrate to film and reduced the interfacial stress resulting in excellent wear resistance when sliding against 6 mm alumina balls [20].

With appropriate tuning of deposition conditions many nanocomposite thin film systems such as a-C, TiC/a-C, TiC/a-C:H, nc-TiC/a-C(Al), YSZ/Au, TiFeN and nc-TiN/a-Si 3 N 4 can have the combination of (1) high plasticity in indentation (2) high toughness in applications and (3) H ~15-20 GPa. Dissipative mechanisms such as grain boundary sliding, which is possible due to this moderate hardness, result in high toughness and resistance to extreme cracking at high contact load. Suppressing dislocation activity to produce harder and super-hard films may be detrimental to their durability under highly loaded conditions, particularly if these films are deposited on softer and more compliant substrates.

Conclusions

The resistance to fatigue fracture by nano-impact testing enables an assessment of the relative toughness of TiFeN films on silicon under dynamic loading. From individual impact depth vs. time graphs it has been shown that film behaviour can be assessed during three distinct regimes as the test progresses: (i) load support (pre-impact) (ii)

resistance to a single (initial) impact and (iii) resistance to subsequent impacts. Film behaviour varies dramatically as a function of impact force and number of impacts.

The resistance to a single (initial) impact was always improved by increasing the films' resistance to plastic deformation (H 3 /E 2 ), as determined by nanoindentation. On repetitive impacting at low forces films with a higher resistance to plastic deformation were much more resistant to the formation of cracks. However, as the impact force in the test increased these films cracked, behaved no better than films with lower hardness and H 3 /E 2 and subsequently suffer delamination over a large area. The optimum mechanical properties for improved dynamic toughness therefore depend on the severity of the impact conditions. Suppressing dislocation activity to produce yet harder films (very high H 3 /E 2 ) is beneficial for low load contacts but can be detrimental to their durability under highly loaded contacts as suitable dissipative mechanisms are not available.

Figures 3 (

 3 Figures 3(a), (b) and (c) after 6 s at 10 mN on TiFe 0.4 Mo 0.4 N 2, TiFe 0.7 N 1.9 and

Figure captions 1 . 2 .

 12 Figure captions

3 . 2 = triangles. 4 . 5 .

 3245 Illustrative impact depth vs. time graphs for the nanocomposite films at 5-50 mN coil force. (a) 5 mN (b) 10 mN (c) 25 mN (d) 50 mN. Key:-TiFe 0.63 N 1.5 = circles; TiFe 0.69 N 1.44 = crosses; TiFe 0.65 N 1.45 = squares; TiFe 0.7 N 1.9 = diamonds; TiFe 0.4 Mo 0.4 N Mean values of (a) Probe depth prior to initial impact (b) Probe depth after first impact (c) Increase in probe depth due to continued impact; i.e. [depth (final) -depth (after first impact)]. (d) Final impact depth. The key for 4 (a-d) is as in figure 3. Variation in measured H 3 /E r 2 (film and substrate composite response) with increasing depth from multi-cycle (load partial unload) nanoindentation to 20 mN. 6. XRD spectra of films.

Figure 1 Figure 6 (

 16 Figure 1 (fig1.tif)

Table 1

 1 summarises the mechanical properties of the films determined by load-partial unload nanoindentation. Values of film elastic modulus independent of substrate influence were obtained by a linear extrapolation to zero-depth. Film hardness varied from 16-24 GPa. With H film and E r,,film known it was possible to determine film-only values of (H/E r, ) film and (H 3 /E r 2 ) film (Table

Table 1 Composition and mechanical properties of TiFeN and TiFeMoN films

 1 In Table1the values of H, E, H/E r and H 3 /E r 2 are all film-only properties determined from load-partial unload nanoindentation tests. Film hardness was determined from the plateau (load-invariant) region and film elastic modulus by extrapolation to zero depth.

	Composition	Ion	pN 2 (Pa) Thickness	H	E r	H/E r	H 3 /E r 2
		Assist/V		(nm)	(GPa)	(GPa)		(GPa)
	TiFe 0.69 N 1.44	280	0.011	1700	23.7	254	0.093	0.206
	TiFe 0.65 N 1.45	300	0.011	1400	22.3	241	0.093	0.191
	TiFe 0.4 Mo 0.4 N 2	-	0.011	~1650	19.8	220	0.090	0.160
	TiFe 0.63 N 1.5	600	0.021	700	20.0	240	0.083	0.139
	TiFe 0.7 N 1.9	600	0.021	920	16.0	192	0.083	0.111

Acknowledgements

Dr O. Crisan is gratefully acknowledged for performing the XRD analysis.


K J Cheetham

A Krier
email: a.krier@lancaster.ac.uk

I Patel

F L Martin

J-S Tzeng

C-J Wu

H-H Lin

Raman scattering in InAs x Sb y P 1-x-y alloys grown by gas source MBE

Keywords: 33.20.Fb, 63.20.-e, 68.55.ag, 78.20.-e, 78.30.-j, 78.30.Fs, 78.66.Fd

The Raman spectra of quaternary InAs x Sb y P 1-x-y epitaxial layers nominally lattice matched to (100) n-type InAs substrates are reported. The Phonon peaks are identified for alloys covering a wide composition range extending into the miscibility gap. Three-mode behaviour was obtained across part of the composition range. InP-like LO and TO phonons were observed over the entire compositional range, but InAs-like LO phonons were only observed at high arsenic concentrations and no InSb-like TO phonons were observed. Disorder related phonon peaks were obtained for alloy compositions within the miscibility gap.

Introduction

The III-V quaternary alloys InAsSbP can conveniently provide access to semiconductors with energy band-gaps within the technologically important mid-infrared (2-5 µm) spectral range [1][2]. The presence of a fourth element also allows lattice-matched growth on InAs substrates. Using this approach it has been possible to produce mid-infrared LEDs [3][4][5], diode lasers [6][7], and photodetectors for detection of gases such as methane [8], carbon monoxide [9] or HF [10] as well as thermophotovoltaic cells [START_REF] Cheetham | Solar Energy Materials and Solar Cells[END_REF]. However, there has been very little fundamental investigation of the InAsSbP material system and particularly for alloys grown by molecular beam epitaxy (MBE). In this work we report on the systematic spectroscopic investigation of InAsSbP alloys grown using gas source MBE. In particular we studied alloys which were lattice matched onto InAs substrates, covering almost the entire composition range and penetrating inside the miscibility gap [START_REF] Wilson | [END_REF].

Raman and infrared spectroscopy are powerful methods for the investigation of semiconductor alloys. For example, lattice vibrations at the characteristic longitudinal optical (LO) and transverse optical (TO) phonon frequencies can convey useful structural information about alloy composition, strain, bonding and the local environment of structure defects and impurities.

It is well known that for zinc blende semiconductors grown on (100) substrates, the TO mode is forbidden in Raman scattering due to selection rules. However, the weakening of selection rules caused by structural disorder or substrate misorientation breaks the symmetry and leads to the appearance of TO modes.

The Raman spectra of InAsSb, InPSb and InPAs III-V ternary alloys have been thoroughly investigated [13][14][15]. However, previous Raman studies on III-V quaternaries have focussed more on gallium containing compounds, such as GaInAsSb [16][17]. Whereas photoluminsence of InAsSbP has been extensively studied [18][19], vibrational studies are rare and limited to a small range of compositions [20][21]. Here we report on the detailed Raman spectroscopy of InAsPSb across a wide composition range.

Experimental procedures

Confidential: not for distribution. Submitted to IOP Publishing for peer review 9 December 2010

The samples for our investigation were grown by gas source MBE on (100) InAs substrates at a growth temperature of 470 °C, using a VG-V80H reactor. The indium was provided by a thermal effusion cell, the antimony by an EPI effusion cell, and the arsenic and phosphorus came from cracked arsine and phosine gasses respectively, as detailed in a previous paper [16]. The fluxes of each of the materials were varied, with the phosphine adjusted from 990 to 300 Torr, to give a wide range of alloy compositions, all approximately lattice matched to InAs. In-situ monitoring of growth was carried out using reflection high energy electron diffraction to observe the condition of the growing epitaxial layer. The resulting layers were all ~1 µm thick and intentionally undoped.

The composition of the samples was determined using a JEOL JXA-8200 electron probe micro-analyser (EPMA) and are given in table 1. The system used for gathering the Raman spectra was a Renishaw InVia Raman spectrometer equipped with a 576x400 pixel Peltier cooled (-70 °C) CCD detector. A spot size of ~1.6 µm 2 was obtained using the 50X objective lens of a Leica DM2500M compound microscope. The spectral resolution obtained was ~1 cm -1 , using a 1200 line/mm diffraction grating and a 50 µm entrance slit. A 100 mW diode laser was used operating at 785 nm (1.58 eV, higher than the band-gap of all the tested InAsSbP samples). The power incident on the sample was ~40 mW after allowing for optical losses. A holographic notch filter was used to block out unwanted reflections. Scans from 100 to 800 cm -1 were taken at several locations on each sample, each with a run of 16 accumulations to reduce the effects of noise. An exposure time of 10 s was used to minimize heating effects with the sample at room temperature, however some local lattice heating is inevitable and results in peak broadening of the spectra.

Sample

Results and discussion

Figure 1 shows the composition plane, along with spinodal curves calculated for different growth temperatures [22]. It can be clearly seen that the four low arsenic samples are therefore grown inside the immiscibility region. Surface roughness can increase the background noise of Raman spectra due to enhanced diffuse scattering. In our case a large number of pits are observed at low arsenic concentrations and are probably due to arsenic vacancies. The optical quality of the sample surface was found to be improved by using an increased concentration of arsenic during growth. X-ray diffraction analysis revealed evidence of spinodal decomposition which has been described in detail in our previous publication [19].

The lattice constant of InAs is 6.0583 Å, and the calculated lattice mismatch of the grown layers was found to be less than 0.4%, with only the samples within the immiscibility region having a lattice constant smaller than that of InAs. Given the samples were all nominally lattice-matched and assumed to be relaxed, the material stress was assumed to not contribute to the Raman spectra at the surface. Similarly, although the InAsSbP/InAs interface is thought to be type II [25], in our case the layers are sufficiently thick so that the substrate plays no part in the Raman spectroscopy. However, the peaks can become broadened due to the stresses induced from lattice mismatch. The optical penetration depth, d opt =1/2α, varies from 66 nm for InAs [26] to 85 nm for InSb 0.32 P 0.68 .

For a mixed alloy AB x C 1-x , a simple criterion for predicting the mode behaviour has been proposed by I. F. Chang [27] to this rule, InAsSb shows one-mode behaviour when the Sb mole fraction is smaller than 0.6 [13]. This concept can be expanded to quaternaries as well. Consequently, we would expect to see three Raman modes, from the InAs-, InSb-and InP-like binaries. Each of the corresponding binary compounds have the same group III element, and the Raman shift is roughly proportional to the inverse of the square root of the atomic mass of the group V element.

The samples grown in the immiscibility region, labelled C1900, C1898, C1899 and C1767, were clearly distinct from the other samples at low energies (fig. 2). They exhibit two principal peaks, labelled A and B, at approximately 138 and 145 cm -1 , which are associated with disorder-activated longitudinal acoustic phonons (DALA), also seen in InAsSb ternary alloys [13], as well as a more prominent InSb-like LO peak. The InAs-like TO peak is forbidden in this geometry (since InAsSbP is a zinc blende structure), but is present in the spectra due to disorder. Several smaller peaks become apparent outside the immiscibility region. The peak at 115 cm -1 is thought to be from the E 1 mode of pure antimony. The A 1g phonon mode of antimony occurs at 154 cm -1 [28], but is only noticeable in one sample as its intensity is too low compared to the DALA peaks. Two further peaks occur at 122 and 129 cm -1 , of unknown origin, but are expected to be disorder related.

The InP-like LO peak becomes more noticeable with increasing phosphorus content, as expected, with an extra DAO peak inside the immiscibility region [29].

The spectra were deconvoluted into Lorentzian curves in order to better resolve peaks which were close to each other. As the forces between atoms depend on their neighbours, in a disordered alloy the phonon frequencies depend on the composition.

Figure 3 shows that both the InP-like TO and LO phonon peaks are observed throughout almost the entire compositional range and exhibited 0.182 and 0.256 cm -1 peak shifts per percent of phosphorus respectively. When the antimony and phosphorus mole fractions approach zero, the LO and TO modes converge to an isolated vibration mode, which is that of an isolated phosphorus atom in InAs, occurring at 310 cm -1 [29]. With increasing phosphorus content a further shoulder appears between the two peaks at 314 cm -1 and is associated with disorder activated optical (DAO) phonons [30]. The Raman signal from samples within the immiscibility region is much broader than those from samples outside the region. Peak D is very weak and is thought to be a second order effect associated with InSb-like 2LO phonons.

Peak C only appeared outside the of immiscibility region, with a very low intensity and was relatively insensitive to composition. As it only appears when the arsenic concentration is >0.6 it is thought to be associated with InAs. A phonon peak at 270 cm -1 was assigned to additive combination TO+TA ( K) in InAs with a nearby 2LA peak by R. Carles [31]. 2 were also included.) The InAslike LO peak was only observed at high arsenic concentrations, with a linear shift of 0.229 cm -1 per percent of arsenic. But, the TO peak was found to more closely fit the parabolic dependence described by equation ( 4). (We were unable to distinguish the TO peak for the second sample from the nearby peak.) Another DAO peak appears as a shoulder between the InAs-like TO and LO peaks. Such a DAO peak has been previously reported in similar materials [32], and was observed here with a greater intensity than even the InAs-like TO for samples C1900 and C1767. There was no peak observed for C1898 (InAs 0.0408 Sb 0.2952 P 0.664 ), we suppose that this is because it becomes indistinguishable from the DAO peak.

ν InAs-TO (x) cm -1 = 211.4 + 14.6x -8.5x 2

The full width at half-maximum (FWHM) of the InAs-like TO peak varied from about 9 cm -1 at very low arsenic concentration to ~20 cm -1 for the arsenic-rich samples. The FWHM for the InAs-like LO peak remained constant at ~4 cm -1 , while the nearby DAO stayed constant at ~6 cm -1 .The FWHM of the InP-like LO peak was found to be approximately constant at ~9 cm -1 , which indicates good crystalline quality, while the nearby DAO was slightly wider at ~15 cm -1 . However, the FWHM of the InP-like TO peak increased to ~20 cm -1 with increasing phosphorus concentration, suggesting increased alloy disorder.

Investigating the antimony component, as shown in figure 4(b), revealed that the Raman shift of the InSb-like LO peak had similar linear dependence on composition as the InP-like peaks.

ν InSb-LO (y) cm -1 = 183 + 20y (0.05 y 0.30) (5) The InAs-like LO peak is more easily observed than the InSb-like LO peak for similar alloy compositions of arsenic and antimony due to the polarization field being larger for higher energy modes, and hence the electron-phonon interaction is enhanced [33]. The second-order 2LO peak from InAs at ~478 cm -1 was not easily observed and was thought to be too weak to distinguish from the background noise.

Conclusion

In summary, quaternary InAs x Sb y P 1-x-y alloys (covering a wide compositional range of 0.040 x 0.795, 0.099 y 0.285) which were grown by MBE nominally lattice matched to InAs were investigated using room temperature Raman spectroscopy. The alloy was confirmed to exhibit three-mode behaviour across part of the composition range. InP-like LO and TO phonons were observed over the entire compositional range, InAs-like LO phonons were only observed at high arsenic concentrations and no InSb-like TO phonons were observed at all. Increased DALA phonon signals were observed for alloy compositions within the miscibility gap. It is evident that Raman spectroscopy can provide detailed and rapid diagnostic information about structural disorder and the onset of phase separation in this technologically important alloy system.

Figure 1 .

 1 Figure 1. Experimental data and the InAsSbP alloy composition plane. Layers from this work are represented by black squares. The solid red line shows lattice matching to InAs.

Figure 2 .

 2 Figure 2. Room temperature Raman spectra of all the samples, with lines added as a guide to the eye showing the peak trends.

Figure 2

 2 Figure2shows the Raman spectra measured from each of the samples. The peak assignment is primarily based on data from table2. With an increase in the mole fraction of arsenic the InAs-like features dominate, as the corresponding modes have increased scattering intensity, and their frequency is closer to that of the end-point binary. The intensities of the peaks give an indication of the mode strength.

  in terms of the individual mass, denoted by m, and the reduced mass µ. Such alloys show one-mode or partial two-mode behaviour if two conditions are satisfied: m B > µ AC and m C > µ AB . Alloys show two-mode behaviour if m B > µ AC and m C < µ AB (m C < m B ). According

Figure 3 .

 3 Figure 3. Raman peak shift with varying phosphorus concentration; the solid lines are linear fits.

Figure 4 .

 4 Figure 4. Raman peak shift with varying concentrations of (a) arsenic and (b) antimony in the alloy. The solid lines are linear and quadratic fits. The circled point is a combination of two indistinguishable peaks.

Figure 4 (

 4 Figure 4(a) shows the dependence of the InAs-like phonon peaks on the arsenic content in the alloy. (The values for pure InAs from table2were also included.) The InAslike LO peak was only observed at high arsenic concentrations, with a linear shift of 0.229 cm -1 per percent of arsenic. But, the TO peak was found to more closely fit the parabolic dependence described by equation (4). (We were unable to distinguish the TO peak for the second sample from the nearby peak.) Another DAO peak appears as a shoulder between the InAs-like TO and LO peaks. Such a DAO peak

ν

  InAs-LO (x) cm -1 = 215.5 + 23x (0.65 x 1.0)

Figure 1 (Figure 2 (Figure 3 (

 123 Figure 1 (Figure1.tif)

Table 1 .

 1 Values of sample compositions determined from EMPA.

		As	Sb	P
	C1898	0.040	0.285	0.675
	C1899	0.092	0.283	0.625
	C1900	0.223	0.260	0.517
	C1767	0.361	0.228	0.411
	C1768	0.560	0.155	0.284
	C1770	0.681	0.099	0.220
	C2054	0.694	0.108	0.199
	C2055	0.730	0.109	0.161

Binary compound TO phonon (cm -1 ) LO phonon (cm -1 )

				Atomic
				weight of
				group V
	InP [21]	306.9	343.6	30.97
	InAs [20]	217.3	238.8	74.92
	InSb [20]	179.8	190.8	121.76

Table 2 .

 2 Phonon frequencies of the binary constituents, plus atomic weights of the corresponding group V atoms.

Acknowledgements

The authors are grateful to the Engineering and Physical Sciences Research Council (EPSRC) for providing a studentship for K. J. Cheetham and to Dr. Frank Martin at the Lancaster University Centre for Biophotonics for providing access to the Raman facilities.