Mechanisms of r.f. plasma nitriding of Ti-6Al-4V alloy

transcript

Materials Science and Engineering, A 167 (1993) 155-164 155

Mechanism of r.f. plasma nitriding of Ti-6A1-4V alloy

A. Raveh NRC-Negev, Division of Chemistry, P.O. Box 9001, Beer-Sheva 84190 (Israel)

(Received November 211, 1992)

Abstract

The objective of the current study was the gradual development of the formation of the nitride layer during inductive r.f. plasma nitriding. The study centers on characterization of refined layers and plasma diagnostics in the vicinity of the sample, and raises critical questions of how the layers and interracial microstructure might affect the near-surface properties. The composition of the plasma near the surface of the sample (plasma layer) was examined by optical emission spectroscopy and mass spectrometry during plasma nitriding and while sputtering the sample after the nitriding process. It was observed that during the nitriding process, the plasma layer contains Ti, NH,, species, N (or/and N + ), H,, species (or/ and H +2). However, when the nitrided sample was exposed to argon plasma, Ti, A1 and NH were observed. It was found that two distinct sublayers, comprising d-TiN and &TiN + e-Ti~N phases, were formed with alloying elements in a segre- gated zone, followed by a solid solution of nitrogen in titanium. The formation of the uppermost sublayer (&TIN phase), containing H, NH, and N, in addition to Ti depleted of AI and V, has a strong effect on the diffusion of nitrogen into a-Ti and on the layer properties. This can be enhanced if hydrogen is present in the nitriding atmosphere and is prevented if hydrogen is replaced by argon. Therefore, the nitrogen content in the layer results in the formation of nitride phases and is accompanied by an improvement in mechanical properties.

1. Introduction

Beyond continuous efforts in developing advance processing methods or new directions in surface modification, the foundations for the appropriate assess- ment of surface layers still remain very challenging. In this context, Ti -6A1-4V ct/fl alloy was mainly investigated after plasma nitriding by nitrogen or by a mixture of nitrogen and hydrogen, and/or argon [1-4]. How- ever, the mechanism of surface nitriding of titanium ct/fl alloy remains an open question. In order to gain more understanding of this issue, emphasis was given to establishing refined scales for layers and local sublayers, chemical and microalloying aspects, and diagnostics of the plasma composition, near the surface of the sample (plasma layer), during plasma nitriding and while sputtering the sample after the nitriding process.

It has been reported [5-13] that the surface nitriding of metals such as steels, titanium, zirconium, silicon and aluminum may be attributed to bombardment of NNH or N + ions [11] and to surface adsorption of N and NH radicals, followed by enrichment of the surface region with nitrogen and diffusion of nitrogen into the material.

Rie and Lampe [14, 15] investigated the nitriding of titanium and of its alloys in a plasma of N 2 + H 2 at temperatures between 600 and 900 °C. They showed that

the nitrided layer is composed of two different layers, a compound layer (CL) and a diffusion layer (DL). The CL was thicker in titanium than in the alloys. The opposite behavior was observed for the DL. Laidani et al. [16] investigated the kinetics of the nitriding of titanium with NH 3 at 40 MHz at 625-850 °C and suggested that the nitride is formed with two activation energies. One is for the formation of the outer layer and the second is for the diffusion process. They also suggested that the composition and the electrical properties of the outer layer accelerated the diffusion of nitrogen into the titanium. Metin and Inal [17] also studied the kinetics of layer growth in ion-nitrided pure titanium at 800-1100 °C. They showed that nitride phases were formed consistent with equilibrium phase diagrams. They calculated the diffusion coefficient D of nitrogen in the Ti-TiN system, and expressed the relations of D~, D,. and D a, where the subscripts a, c and d refer to the nitride phase.

Recently, an intensive investigation has been initiated of plasma nitriding of a/fl titanium alloys [4, 18-20] at 440-540 °C, as well as the nitriding of pure titanium and its alloys at 800-1000 °C. Generally, the program was centered on layer characterization and ways of refining experimental findings aimed at sub- stantiating the role of processing variables and their effects. In order to assist in achieving these objectives,

Elsevier Sequoia

156 A. Raveh / Plasma nitriding of Ti-6AI-4V

various techniques were used such as optical emission spectroscopy, Auger electron spectroscopy, X-ray dif- fraction (XRD), scanning and transmission electron microscopies (SEM, TEM). Results on Ti-6A1-4V and Ti-8A1-1V-1Mo in the lower temperature regime under certain conditions, such as long time of pretreatment by argon + H 2 plasma and substrate bias, have indicated that different mechanisms are involved at low and high temperatures, as manifested by distinct effects of the plasma together with different features of the sublayers [14, 18].

Accordingly, the current investigation of plasma nitrided a/fl titanium alloy was motivated by the wish to correlate plasma processing, the microstructure and mechanical properties of the inductive r.f. plasma nitrided alloy at low temperature (440-540 °C), and to develop some understanding regarding the mechanism of layer formation.

2. Experimental procedures

Intensive experimentation was carried out on Ti-6AI-4V, containing (in weight per cent) 6.0% A1, 4.0% V, 0.03% Cu, 0.01% Cr, 0.32% Fe and Ti (balance). The plasma nitriding set-up is shown in Fig. l(a). Two r.f. generators were used: a 50 MHz generator operating at 1.1 kVA, or a 13.56 MHz generator operating at 300 W. Substrates were positioned in one of three locations along the reactor: position H, upstream, where fresh reactant is introduced; position G, central, at the center of the coil; position F, down- stream. The dimensions of the samples for nitriding were 40 × 10 x 5 mm 3. The samples were lapped down to 0.25 /~m with diamond suspension then degreased ultrasonically in acetone and dried in air.

Prior to the nitriding process, the samples were acti- vated by surface sputter cleaning in an argon and hydrogen (2:1 ratio) plasma for 60 min. The nitriding was performed by N2, N 2 : H e = 4:1, N 2 : H 2 : Ar = 8 : 1 : 1 or N 2 : A r = 4 : I plasmas. The total flow rate varied beween 50 and 150 standard cm 3 min -~ with a bias voltage which varied between 0 and - 300 V d.c. for a nitriding time ranging between 10 min and 10 h. The temperatures of the samples were measured directly by a thermocouple introduced into the bulk of the samples, and by an optical pyrometer. The measured temperatures of the samples during the nitriding process were 440___20 °C at positions H or F, and 470 _ 20 °C at position G. The maximum bias voltage ( - 300 V d.c.) increases the temperature of the samples up to 490 _ 20 °C at H or F, or up to 540 +_ 20 °C at G.

In order to investigate the concentration of excited and ionized particles of the gas mixture in the plasma layer (PL, 0.2 cm from the surface of the substrate)

GAS INLET SYSTEH

Av H 2 N 2 ~ I

R.E GENE

PiSSURE GAGE

TEMPERATURE GAGE I

VACUUM SYSTEM ]

TEMPERATURE GAGE ] I

*I I ', VAGUU Y TE. ]

GENERATOR "

- - - o_r

GAS INLET SYSTEM

(b) A, N2 "z

Fig. 1. Plasma nitriding set-up (a) and diagram of plasma diagnostic system (b). OES optical emission spectroscope, QMS quadrupole mass spectroscope.

prior, during and after the nitriding process, a microwave plasma was also utilized. The N2 :H2 mixture of the microwave plasma and the mixture of sputtered and etched particles were examined by quadrupole mass spectrometry (QMS) and optical emission spectroscopy (OES) in the system layout as illustrated in Fig. l(b).

The emission of spectral lines of sputtered and etched atoms such as titanium, aluminum and vanadium were measured in the PL, focused on the entrance slit of a computerized optical emission system. The spectra of Ar, H, N 2 and NH were also monitored as a function of pressure (1-9 Torr), power input (50-200 W) and bias voltage (0 to - 300 V d.c.). Mass spectra of the above species were measured in the PL by computerized QMS; normalized relative concentrations of the measured species were evaluated. The results thus obtained shed light on the surface composition and processes taking place in the immedi- ate vicinity of the sample.

A. Raveh / Plasma nitriding of Ti-6Al-4V 157

After the nitriding process, the layer structure, properties and microstructure of the surfaces were characterized by various techniques including XRD, Auger electron spectroscopy, SEM and TEM, as described in detail elsewhere [21-23]. The whole program has been extended to comparative studies between plasma and thermal nitriding layers (without plasma).

3. Experimental results

NH~ species increases by applying a bias of - 3 0 0 V d.c., while the concentration of NH~- species remains unchanged. Increasing the pressure up to 5 Torr also increases the relative concentration of NH~- and NH~- species. At a higher pressure the NHJ- concentration

3.1. Plasma characteristics The effect of hydrogen during plasma nitriding by an

N: + H 2 mixture was studied at 1-9 Torr with mixtures of N 2"H 2=4:1 and N 2:H 2=1:4. Table 1 presents mass spectrometry data for a characteristic spectrum obtained at 1 Torr and shows the distinct effect of the "g plasma. Mainly masses 16 and 17 increased after plasma activation, compared with the gas mixture with- ~ out plasma. The appearance of masses 16 (NH~) and 17 (NH~) indicates that the reaction between N 2 and H2 has occurred and that this reaction is enhanced by the negative bias. The masses 2 (H~-) and 28 (N~-) show similar intensities (approximately 4000 a.u.) under different plasma conditions at a pressure greater than 2 Torr, and only masses 16 and 17 are meaningful. The variation of the intensities of masses 16 and 17 is shown in Fig. 2. It was observed that with the N 2 : H 2 = 1: 4 plasma, the intensities of masses 16 and 17 increased up to five times compared with the 100 N 2 : H 2 = 4" 1 plasma. This also indicates that a higher content of NH~- and NH~- species is formed with increasing hydrogen concentration in the N 2 + H 2 " ~ 1 0

plasma. The effect of the pressure on the relative concentration (I/ZI) of the different species with grounded .~ (0 V d.c.) and biased ( - 300 V d.c.) substrate is shown in Fig. 3. As seen from this figure, the concentration of Z 1

TABLE 1. Spectra of the masses observed in the N 2 : H 2 = 4: 1 0.1 mixture, pressure 1 Torr, power 150 W 100

Intensity (arbitrary units) m/e Identification

a b c d

2 H f 676 598 585 571 14 N ÷ 21 24 25 25 16 NH~ 0 1 2 4 17 NH~ 0 10 12 19 18 NHJ-, H20 + 61 63 68 61 28 N 2 +

2674 2715 2677 2765

a, Gas composition without plasma; b, plasma without sample; c, plasma with grounded sample; d, plasma with sample biased at - 300 V d.c.

t(A) ' ' = 1:4 d N2:H 2

a b c d

I---T1, ,

i( ' ,

B) N2:tt 2 = 4:1

I a b c d a

a b c d

o m / e 14 16 17 18

+ N H + + + ion identification N + NH 2 3 H O ,NH 4

Fig. 2. Mass spectra during nitriding, (A) N 2 : H 2 = l : 4 , (B) N z : H 2 = 4:1, 5 Torr: (a) gas mixture without plasma; (b) plasma without sample; (c) plasma with grounded sample; (d) plasma with sample biased at - 300 V d.c.

I I I I 1 ~ l I ~ ~ I I I [ I l

W (a) , . - - - - N +

/ Nil* 2

[ I I I

- - - I q . ~ - _ _ _ J _ - [ - _ _ _ _ _ .

H ÷ (b) 2

"'7.. . . . . . . . . . . . . . . . . . . . N + = I 0 - ' "

~ ~ . . . . . ..-'" NH 3 := / I ~ - [ +

o.1 i i i i

0 2 4 6 8 10 Gas Pressure (Torr)

Fig. 3. Mass spectra of N 2 :H 2 = 4:1 plasma vs. gas pressure: (a) grounded sample; (b) bias substrate of - 300 V d.c.

158 A. Raveh / Plasma nitriding of Ti-6Al-4V

decreases for a grounded substrate, whereas it reaches a constant value at a bias of - 300 V d.c. The increase in concentration of NH~- and NH~- species is accompanied by an increase in concentration of N ÷ and a small reduction in H~- concentration.

Avni and Spalvins [11] and Hudis [5] showed that under similar conditions the concentration of ion + atomic species reached a maximum value in the N2 : H2 = 1:1 plasma compared with N 2 + argon plasma and with N 2 plasma. They also found that the addition of H2 to the N 2 + argon plasma increases the relative concentration of N, N 2, NH2, NH3, etc. of ion, excited or neutral species. The increase in the relative concentration of the energetic species i n the plasma increases the gas decomposition rate and the formation of ion, excited or neutral species [24, 25].

OES during sputtering revealed that only titanium was sputtered from the untreated sample, while spectra of aluminum, NH, Ha and H~ were identified from nitrided samples. The concentrations of titanium and aluminum were higher at a high bias ( - 300 V), while that of NH, H a and H a reached a maximum value at about - 1 5 0 V (Fig. 4(a)). This is probably due to

recombination processes at biases greater than - 1 5 0 V. Furthermore, during nitriding NH, N 2 and titanium were detected, indicating that the surface of the sample is sputtered during the process. The etched and sputteced titanium or the already formed nitride can react in the plasma layer (homogeneous reaction) and deposits as a TiN layer. It was also observed that titanium, aluminum and NH species were detected in the depth of the layer, while NH was only detected down to a depth of 0.5/~m, as shown in Fig. 4(b).

3.2. Process-dependent structure and property The samples nitrided in various nitriding atmo-

spheres show in addition to a titanium-rich solid solution a'-(Ti, N), two additional phases (Fig. 5). These include the e-Ti2N and the 6-TIN phases. It is observed that the dominant peaks of films nitrided at region G (Fig. 5(c)) belong to the &TiN phase whereas these peaks are not detected in films at F (Fig. 5(b)). How- ever, titanium hydride or intermetallic compounds of the alloying elements as their hydrides or nitrides are definitely absent. In order to determine the plasma pro-

~ 1 0 0

( a ) , T~

" Ar • A1

100 200 300 400 Bias (-Volt)

~ 40 \ N i t (1)

"= 2 0 - )

0 - , - , " , - , -

0.15 0.30 0.45 0.60 0.75

Depth (~un)

Fig. 4. Optical emission in the plasma layer during sputtering in argon plasma, Ti-6A1-4V sample nitrided at G in N2:H2 = 4:1 plasma: (a) as a function of bias; (b) as a function of the nitrided depth, sputtering at - 3 0 0 V d.c. Ti 3998 A; AI(1) 3961 A; Al(2) = 3944 A; NH(1) 3360 A; NH(2)3370 A.

(b) , 8-Ti -Ti2N

90 80 70 60 50 40 30

Dif f rac t ion A n g l e ( 2 0 )

Fig. 5. X-ray diffractograms of nitrided Ti-6AI-4V, N 2 :H a = 4:1, frequency 0.5 MHz, pressure 5 Torr, nitriding time 5 h: (a) untreated; (b) nitrided at F; (c) nitrided at G.

A. Raveh / Plasma nitriding o f T i - 6 A I - 4 V

TABLE 2. Effect of processing parameters on the structure and properties of plasma nitrided a/f l Ti alloys

Structure and property Processing parameter Reference

d-phase G, bias substrate, frequency, N 2 + H 2 p l a s m a 18, 20 e + d phase G, grounded substrate, or F, bias substrate 4, 19, 20 a' or a' + e F, grounded substrate, structure of the alloy 4, 22 Phase composition and orientation Gas composition, frequency 19, 23 Diffusion rate (thickness) Surface activation, bias substrate, frequency, gas composition 20, 23 Surface microhardness Phase composition, concentration 18 Erosion resistance Phase composition, d-phase 4, 26 Fatigue resistance Phase composition, a' or a' + e 26

cessing parameters which affect the microstructure and properties of the nitrided alloy, some results are correlated in Table 2. The table shows that the location of the sample along the reactor (G or F) is a controlling factor in the formation of the nitride phases. In addition, gas composition affects the orientation and relative concentration of the nitride phases. For example, by nitriding in an N 2 + H 2 plasma a randomly oriented d-phase is obtained together with an e-phase which is strongly oriented in the (Q02) direction [18, 21]. Changing the nitriding atmosphere to N2+argon 60 causes a texture reversal, i.e. the b-phase is strongly ~" 5 o

oriented in the (200) direction and the e-phase is almost randomly oriented. ~ 40

Nitriding at position G in a plasma of N 2 + H 2 ~ 30 mainly causes the formation of the d-phase together ~ 20 with the e-phase, with a very small amount of the a '- phase (Fig. 5(c)). The addition of argon to the N z + H 2 0 plasma results in an inversion of the phase composition, the d-phase being in the lowest concentration [4, (a) 22]. In a plasma of argon + N 2 the three phases are 60 obtained in equal amounts. However, nitriding at posi- g 50 tion F mainly results in formation of the a '-phase with

40 small amounts of the e-phase (Fig. 5(b)). The composi- o =

",= 3 0 tion of the gas mixture only has a small effect on the phase composition obtained at position F [23]. ~ 2o

After plasma nitriding at position G, the layer was ~ 10 composed as follows: 0

(a) (near surface) a diffusion sublayer (DL) of about 2 -20 /~m consisting of mainly a,-(Ti, N) or sometimes (b) with the e-phase as well;

(b) more of a far field affected zone, a defined com- ~. pound layer (CL) of nitrides was identified with two ~ 40

sublayers, (i) about 2 -8 /~m of lamellar microstructure ~ 3 o

containing e + b phases, and (ii) about 0.5/~m in scale =o a0 of a columnar d-TiN phase, g 1 o

Generally, alloying elemental segregation occurs in the interracial domain confined between two sublayers 0 0 [19, 23]. As shown and discussed [4], the complete (e) structure is only obtained when nitriding at position G, while the CL and segregate zone do not form at position F. However, nitriding a biased substrate at position

F can produce a similar structure to that formed at position G for a grounded substrate. Still, the formation rate and the formation of the nitride phases, e- Ti2N + b-TiN, are less at position F compared with position G.

The effect of the location of the nitrided sample along the reactor on the uppermost surface composition is shown in Figs. 6 and 7, enabling some connec-

I I I ~ I - - - - . ~ i

- - - 1 . . . . I - - - - T - - - - F - - - - - I V

0.2 0.4 0.6 0.8 1

Depth (gm)

I I I I i

. . . . . . . . . . - - . . . . Y i

; - = - - L ] _ _

A1 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . V

. . . . ! . . . . . 1 - - - - " l - - - - - I . . . . l

0.2 0.4 0.6 0.8 1

Depth (p.m) 6 0 , , , , ,

5 0 ~-- . . . . . . . . . . . . Ti 1

/ . . . . . . . . . . . . . .

-:"-:"" . . . . . . . . . . . . "" " - N

• .... AI - - - - t - - ' - " - - I - - - - - r - - - - - r . . . . . I V

0 . 2 0 . 4 0 . 6 0 . 8 1

Depth (gm)

Fig. 6. AES depth profile of plasma nitrided Ti-6AI-4V, N2:H 2 = 4:1 plasma, 13.56 MHz: (a) position H; (b) position G; (c) position E

tions to be made between processing variables and product properties. For the sample nitrided at region H, it was observed that the Ti:N ratio increases into the layer. The aluminum concentration in the surface layer is lower (3-4 at.%) than in the bulk alloy (10 at.%). Similar behavior is observed in the sample nitrided at position G. In the G sample, the aluminum

4 . 2 4 8

4 . 2 4 4

4 . 2 4 0

8 • ~ 4 . 2 3 6

4 . 2 3 2

== 4 . 9 4 4

i 4.940 e~ o • ~ 4 . 9 3 6

4 . 9 3 2

2 . 9 5 0

2 . 9 3 9

2 . 9 2 8

• ~ 2.91 6 ,1

2 . 9 0 5

I I I I I

.0 / f

2 0 30 4 0

/ / "~,

30 4 0 5 0

" ' " 0 ° ' " /

i / 6 " (c)

I I I I I

1 0 2 0 3 0 4 0 5 0

Concentration ( v o l . % )

3 .041

*< 3 . 0 4 0 o t.

3 . 0 3 9 i

3 . 0 3 8 o

3 . 0 3 7 ~

3 . 0 3 6 6 0

4 . 6 8 6

4.681 o

4 . 6 7 5

4 . 6 7 0 ~

,.-.1 4 . 6 6 5

Fig. 7. Lattice parameters of nitride phases vs. hydrogen concentration in N 2 + H 2 mixture, 0.5 MHz: (a) b-TiN; (b) e-Ti2N; (c) a-(Ti, N).

and vanadium concentrations were lower compared with those observed at H. In contrast, the opposite behavior was observed in the sample nitrided at F: the aluminum is enriched at the surface up to 25 at.%. In this region, the Ti:N ratio is about 1:1. As the depth into the layer grown at position F increases, the aluminum concentration decreases and the Ti:N ratio increases, whereas the vanadium content is constant. No such behavior is obtained after nitriding for more than 10 h.

The depth profiles of titanium, nitrogen, aluminum and vanadium in thick layers nitrided for 10 h show that both the aluminum and vanadium concentrations decrease in the nitrided layer and increase at the boun- dary between the nitrided layer and the bulk. The composition of the intermediate region, between the CL and DL, is discussed elsewhere [19].

Figure 7 shows the lattice parameters of the nitride phases as a function of the hydrogen concentration in the N 2 + H 2 gas mixture. The lattice parameters of the e-phase reach a maximum value at 10 vol.% H2, while the lattice parameter of the 6-phase increases continu- ously with increasing hydrogen concentration in the N2 + H2 plasma. For the e- and 6-phases the lattice parameters are higher in the N 2 + H 2 plasma compared with the N 2 + argon or N 2 +H2 + argon plasmas (see Table 2). The lattice parameter (a=4.245 A) of the 6-phase obtained in the N 2 : H 2 = 4:1 plasma indicates overstoichiometry or distortion of the lattice [27]. The increase in lattice parameter of the 6-phase is accompanied by an increase in the lattice parameters a and c of a-(Ti N) (Fig. 7(c)). It was also observed, based on the lattice parameters of pure titanium and of untreated Ti -6AI-4V (Table 3), that adding hydrogen to the plasma depleted the nitrided zone from at least one of its components, aluminum or vanadium. Therefore, it is suggested that hydrogen in the plasma promotes a systematic increase in the lattice parameters toward

TABLE 3. Lattice parameters" of identified phases as a function of gas feed composition: 1, 100 vol.% 2, N 2: H 2 = 80: 20; 3. N2 :HE :Ar = 80:10:10; 4, N2 :Ar = 80:20

Gas b-TiN e-Ti2N a'-Ti composition

a a c (A) (A) (A)

c/a a c c/a (A) (A)

1 N z 4.2328 4.9340 3.0369 2 N 2 + H 2 4.2419 4.9442 3.0407 3 N 2 +H2-Ar 4.2408 4.9421 3.0350 4 N2+Ar 4.2392 4.8143 3.1318

Ti b Ti-6A1-4V

0.6150 2.9073 4.6763 1.6085 0.6143 2.9228 4.6820 1.6019 0.6141 2.9217 4.6747 1.6000 0.6505 2.9229 4.6720 1.5985

2.9510 4.6855 1.5878 2.9239 4.6710 1.5975

aAverage error + 0.0005 A. bData from Pearson [28].

A. Raveh / Plasma nitriding of Ti-6A I-4V 161

pure titanium (see Fig. 7(c) and Table 3), and the solubility of nitrogen in the titanium is due to a reaction between active nitrogen species and the titanium depleted of aluminum and vanadium. In turn, the penetration of nitrogen into the a-Ti is assisted and enhanced by hydrogen if present in the nitriding atmosphere.

Furthermore, XRD analysis shows that nitriding with N 2 + argon o r N 2 + H 2 + argon causes non-systematic changes in the lattice parameters of e-Ti2N and d-TiN phases, compared with the N 2 + H 2 plasma. The addition of argon to the N 2 + H 2 plasma also reduces the solubility of nitrogen in a-Ti and consequently the layer formation is decreased as shown and discussed elsewhere [26]. In addition, argon in the range 0.1%-0.2% is observed in the nitriding surface by energy-dispersive analysis (EDS) performed in the SEM system. Argon incorporation into the nitriding surface can cause stresses or microstrain in the layers, similar to that observed by Bland et al. [29]. Therefore, it is suggested that the penetration of nitrogen into the a-Ti is prevented by argon, if present in the nitriding atmosphere.

Penetration of nitrogen into titanium a/fl alloys has been shown to produce a hardened surface [30] and improved fatigue life [26]. In Ti -6AI-4V it was found that this improvement can be correlated with the hardening by d-TiN precipitates and by compressive elastic stress in the a-Tix-N~ (lattice expansion) [31]. Improvement in the layer properties was found after the formation of e-Ti2N [26].

Figure 8 presents the erosion reduction X and the microhardness H~ of the nitrided layer as a function of the nitrogen content CN at a depth of 0.1/~m. It can be seen that ;t and Hv are directly related to CN. This indi-

£ ¢J

,'C) . - /

4'1" ALl t

0 / 0 /

5 t t 1

25 30 35 40 45 50 55 Nitrogen Content (at.%)

2.s a¢ ¢,-, Q

Fig. 8. Surface microhardness and erosion resistance (4 = 5W1/ 5 W 2, where 5 W~ and 6 W 2 are the weight losses after blasting of the nitr ided and reference samples respectively) as a function of nitrogen content C N in the layer at a depth of 0.1/~m.

cates that high C N, namely the formation of d-TiN or e + d phases, is accompanied by surface hardening and an improvement in resistance to erosion. However, the d-TiN phase may be brittle and reduce the surface fatigue resistance [26]. Therefore, knowledge of the plasma parameters such as the location of the sample in the reactor, feed gas composition, and the structure-proper ty relationship allows us to "tailor-make" nitrided surfaces for different mechanical applications.

4. Discussion

As expected, the thickness of the nitrided layer, the formation of nitrided phases and the properties are dependent on the plasma processing conditions. Parameters such as the exact position in the reactor, generator frequency, bias voltage and the gas feed composition, and other related variables, are also discussed elsewhere [ 18, 19]. At position G, for example, a higher energy is delivered to the plasma than at positions H or E This causes greater decomposition and ionization of the gases and the formation of a higher concentration of energetic particles, resulting in greater penetration of nitrogen into the surface. At high frequencies, an additional increase in the concentration and density of energetic species can be achieved. This is due to a higher population of electrons in the energetic tail of the electron energy distribution function (EEDF) [32], which in turn results in a substantially more efficient plasma reaction rate at high frequencies. Higher plasma reaction rates were also deduced from the thicker compound and diffusion layer, indicating higher nitrogen diffusion. A similar effect of higher deposition rate as a function of frequency has been observed for deposition processes, as described for some carbon films [33].

It was also observed elsewhere [34] that the emission intensities of N, H a and H e in the N 2 or NH 3 plasma follow the changes in the EEDF at lower (r.f.) and higher (microwave) frequencies, in terms of average energy introduced by an electron per unit density of the gas O/P [32]. The value O/P decreases when the frequency increases owing to a higher electron density ne at a given pressure and a fixed observed power. The increase in n e and EEDF favoring elevated electron energy compensates for this drop in ®/P and enhances the total emission intensity. Therefore, more energetic states are excited and the ionization is enhanced, and the proportion of higher energetic species with a higher excitation potential increases (see for example ref. 35). As a consequence, the processes of deposition and surface modification are more efficient at higher frequency discharges [36]. The substrate bias causes higher energetic species in the plasma-surface layer,

and a higher nitrogen content of nitride phases compared with that obtained for grounded substrates. It is noted that hydrogen in the gas mixture increases the nitrogen content in the nitrided layer and enhances nitride formation, while argon inhibits and reduces their formation [21]. This can be explained by the higher concentration of nitrogen-hydrogen molecular ions in the N 2 : H 2 = 4:1 plasma. Furthermore, the gas feed composition has an additional effect on the crystallographic orientation. For example, the addition of argon to N 2 + H 2 plasma changes the (002) e-phase orientation to a random structure [19]. However, the addition of H 2 to the N 2 plasma increases the (002) oriented e-phase and forms a random 6-phase.

The concentration of elements of the alloy near the surface and the structure of the alloy have an important influence on the diffusion of nitrogen into the bulk. This is in addition to the competitive effects of stabiliz- ing the a-phase by nitrogen and the r-phase by hydrogen [37]. Therefore, information on plasma-surface interactions which enhance and modify the surface structure and composition enables nitrides to be pro- duced instead of hydrides. For example, pretreatment by argon + H 2 plasma enhances the depletion of aluminum at position G compared with that observed at positions H or F. The aluminum depletion near the surface can promote the initiation of titanium nitride formation. Furthermore, plasma containing hydrogen promotes the surface activation by diffusion of NH, to a maximum depth of 0.5 k~m, as revealed by OES analysis.

The basic differences between plasma nitriding and thermal treatment result from the energetic species in the plasma phase. In the thermal process, nitriding proceeds by the diffusion of nitrogen molecules. However, in plasma nitriding, the reaction between the plasma and the surface proceeds by nitrogen-hydrogen molecular ions, whereas the chemical and energetic state of the ions and their concentration can vary as a function of plasma parameters. Increasing the concentration of energetic species in the plasma by increasing the frequency, sample position and/or bias voltage can affect the plasma process, while this cannot be performed by conventional nitriding. Nitriding at high temperatures (850-2000 °C) forms a layer structure of 6-, e- and a'-phases (see for example ref. 38). The transition from the a'-phase to the e-phase to the 0- phase occurs as the nitrogen content increases. An equilibrium state is achieved by the formation of each phase as a function of temperature and nitriding time, consistent with the equilibrium phase diagram. How- ever, our results of samples nitrided at position G and low temperature, after prolonged surface pretreatment by argon + H2 plasma, indicate that the plasma has a significant effect at low temperatures, while increasing

the temperature increases the thermal processes compared with the plasma effect.

To conclude, the results suggest that the reaction between nitrogen and titanium is controlled by the concentration and density of energetic species in the plasma, surface structure composition, and chemical element composition in the uppermost surface. The nitriding of titanium alloys with plasma of nitrogen, hydrogen and argon occurs via several parallel paths.

(a) Interaction between the plasma and surfaces of Ti-6A1-4V causes sputtering and etching of titanium from the metal and penetration of energetic species such as N, NH, H and Ar. Pretreatment of the metal by argon + H 2 plasma and/or nitriding by N 2 + H 2 + argon plasma activates, enhances the surface chemical reaction [20], and affects the electric properties of the uppermost layer [16]. The new chemical and electrical properties of the top layer enhance the diffusion of nitrogen into titanium, causing higher nitride phase formation.

(b) The differences between the titanium alloys are governed by the nitrogen and hydrogen diffusion rates into a-Ti and/or fl-Ti [37]. Nitrogen stabilizes the a- phase while hydrogen sabilizes the r-phase [1, 37].

(c) Hydrogen and argon have opposite effects on nitride phase formation; hydrogen in the plasma causes a systematic change in the lattice parameters toward increasing the unit cell volume [18]. Thus, hydrogen enhances the purity of the titanium during the formation of TiN, while the formation of TiN is inhibited by the addition of argon to the plasma.

(d) Nitriding at position G produces a negative segregation of alloying elements which do not form nitride compounds, such as aluminum and vanadium: segregate zones were observed between the CL and DL. However, nitriding at position F causes positive segregation of aluminum to the surface of the alloy and inhibits the formation of the nitrided layer.

As the nitriding occurs along several paths, the formation of the nitrided region is strongly dependent on the plasma conditions. For example, high energetic species at position G compared with F result in a higher nitrogen content at the nitrided surface. Similar behavior is obtained by a substrate bias or with a higher frequency. The phase formation, crystallographic orientation, composition of the layer and its thickness are dependent on several steps: (i) the formation of TiN as a result of the reaction between active nitrogen species and titanium; (ii) depletion of the alloy of aluminum and vanadium and formation of a compound layer of e-Ti2N + 6-TIN; (iii) solution of nitrogen in the a-phase and formation of the a'-(Ti, N) phase.

The addition of hydrogen to the nitriding plasma enhances step (ii), while the addition of argon reduces the solubility of nitrogen in titanium and promotes the

A. Raveh / Plasma nitriding of Ti-6AI-4V 163

formation of the e-phase. Moreover, the spectroscopic observation of a titanium peak in the plasma layer sug- gests the formation of TiN in the plasma layer (homogeneous reaction) and its deposition onto the surface. Therefore, a thin layer of 6-TIN with a fine columnar structure is formed. The sputtering and redeposition of the already existing titanium nitride examined by SEM reveal that the maximum thickness of this layer is 0.5 /~m, even for long nitriding processes, i.e. about 10 h.

Nitriding without plasma at a temperature of 500_+10°C (24-48h) causes the dissolution of nitrogen in the a-phase and formation of an a'-(Ti, N) phase, similar to that observed at step (iii) in the plasma nitriding process. Nitriding at position G exposes the alloy to a high concentration of high energy nitrogen which enhances steps (i) and (ii). However, at position F and with thermal nitriding, energetic nitrogen is present at lower concentrations and cannot form a compound layer.

5. Summary and conclusions

Inductive r.f. plasma nitriding and the formation of nitride phases at low temperature after surface activation were studied. The formation of the sublayers is strongly dependent on processing conditions: sample location in the reactor and the presence of hydrogen and/or argon in the nitrogen plasma. It was observed that the structure-property relationship of the nitrided alloy is affected by the combination of a few factors: phase content, orientation, surface composition, and residual stresses. Thus the following is concluded.

(a) Formation of the nitride phases at low temperature is strongly dependent on the pretreatment for surface activation, concentration of the alloying elements in the near-surface region and structure of the nitrided alloy.

(b) The formation of the nitrided layer proceeds by several simultaneous steps: sputtering and etching of titanium from the surface, reaction between titanium and nitrogen in the gas phase and deposition of TiN, and diffusion of H, NH, and N into the treated surface.

(c) The reaction between active nitrogen species and titanium depleted of aluminum and vanadium forms a compound layer of e- and 6-phases, and has a strong (002) effect on structure and properties.

(d) The structure and properties of nitrided alloys are strongly dependent on plasma conditions. Hydro- gen promotes the formation of e + 6 phases, while argon promotes the formation of an e-phase.

(e) The current study, where refined microstruc- tural characterization of layers and local mechanical evaluation are combined, promises some progress in establishing a nitriding mechanism for titanium a/f l

alloys. This might assist in evaluation of layer perform- ance and processing optimization.

Acknowledgments

A.R. wishes to thank Professors A. Grill and R. Avni for encouragement, helpful and fruitful discussions during the course of this work, Drs. G. Kimmel and P. L. Hansen (Technical University of Denmark) for fruitful discussions on the X-ray and TEM results, N. Mayo, A. Amar, M. Kupiec and D. Rosen for assis- tance in the experimental work, and Z. Barkai for the SEM observations.

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Mechanisms of r.f. plasma nitriding of Ti-6Al-4V alloy

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