(X=Mo, Nb, V, Zr) with Ho.T Gazes. Example of Nitriding.·

TITANIUM'99: SCIENCE AND TECHNOLOGY

A Method for Scanning the Reaction Behaviour of Ti-X alloys (X=Mo, Nb, V, Zr) with Ho.t Gazes. Example of Nitriding.·

A. GUILLOU, J.P. BARS, D. ANSEL, J. DEBUIGNE ..

INSA Rennes, GRCM,.20 Av. des Buttes de Coesmes, F-35043, Rennes Cedex, France.

ANNOTATION

The study of the reaction behavior of binary or more complicated alloys with hot gases is difficult and time consuming. The study of reactions of hot gases with diffusion couples is a method for direct scanning the reaction behavior of alloys of various compositions. Sufficient initial knowledge of diffusion properties allows very valuable results to be obtained. This method is applied to the study of the reaction of the Ti-Mo, Ti-Nb, Ti-V and Ti-Zr beta solid-solutions with nitrogen.

Keys words : Nitriding, Titanium alloys, Beta-stabilizer elements, Diffusion

1. INTRODUCTION

The study of phases grown during diffusion in interdiffusion couples is a well known method for assessing binary phase diagrams. Fast and valuable results can be obtained if: i) the hypothesis of local equilibrium applies, ii) the kinetics of diffusion and growth allow the fonnation of all the phases thennodynamically stable at the interdiffusion temperature. In the reverse, the knowledge of binary phase diagrams allows, for example, the analysis of the oxidation of pure metals and also of dilute alloys to some extent. When systems are no more binary or quasi binary, the study of solid state diffusion and reaction of alloys with hot gases can become very complicated and time consuming. This is a stumbling block as far as the research goal to find an.alloy composition meeting well defined requirements for industrial applications. To overcome these time consuming investigations, we propose an experimental method evolved from our solid state interdiffusion researches : the reaction of hot gases on a cross-section of diffusion couples previously interdiffused. In that way the characteristics of the gas-alloy reactions can be obtained simultaneously over a large composition range. This gives a panoramic view in composition for the reaction of hot gases with alloys. Similar experiments have been carried out by various authors for multicomponent solid state diffusion studies and phase diagram detenninations. In each :case the results · can be graphically represented as diffusion paths in isothermal sections of multicomponent phase diagrams [ 1,2]. Our method is applied to the investigation of the nitriding of binary alloys, at temperatures for which a single solid solution is observed for the whole composition range. As examples we present the nitriding ofTi-X (X = V, Mo, Nb, Zr) beta solid solutions at high temperatures. Studies of the reaction of nitrogen with diffusion couples between other titanium alloys are under development in our laboratory. Other reacting gases can also be used. These studies imply preliminary quantitative interdiffusion detenninations between Ti and X [3,4].

2. EXPERIMENT AL PROCEDURES

Diffusion couples are made of two pellets cut from metals ingots. Faces to be in contact are carefully polished to obtain flat surfaces. The properly diffusion stage is preceded by a short welding stage called prediffusion while the two pellets are mechanically maintained in contact. Only one fast thennal cycle around the a/P transus of titanium is necessary to have a good prediffusion. The diffusion in P phase that gives a large concentration gradient zone is carried out at high temperature in a r.f. furnace (I 0 kHz) in purified argon. The experimental devices have already been described earlier [5]. The diffusion time has been chosen in order to obtain a concentration gradient zone about 4 mm width. Couples are then cut perpendicularly to the diffusion interface and polished. The one part allows the detennination of the diffusion profiles by EPMA and the other is used for the reactions studies with hot gases. Since the Ka N and L! Ti emission lines overlap, we use a deconvolution software. The EPMA determinations are accurate to within 0.5 at.% for the metals and I at.% for N. The

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interdiffusion coefficients are calculated by the Den Broeder method taking into account the molar volume variation. At low concentration the diffusion coefficients are obtained by Hall's method. The other parts of the couples are nitrided in nitrogen (quality N60 supplied by Air Liquide) in a r.f. furnace (500 kHz).

3. INTERDIFFUSION RESULTS

Typical diffusion profiles are given on figure I. The shape of the concentration-penetration curve in the diffusion couple shows the variation of the diffusion coefficient with composition : For the beta stabilizer elements (Mo, Nb; Y), these curves show a strong variation of the diffusion coefficients with respect to concentration and specially for the Ti-Mo system.

100 ..------.---, 100 ..------~~ 1600 ·c. B1oo s 1600 ·c - 14400 s BO . :c BO ~ z '#- :: 60 .!!!. ! 60 c .§ 0 i! c 40 ~ 40 8 c 5c 8 8 20 20

0 200 400 600 BOO 1000 1200 0 200 400 600 800 1000 1200 1400 1600 penetration (µm) penetration (µm)

100 ------100 ------·--- 1 1550 ·c -21600 s 1450 "C • 3BBO s 80 80 ;G'

*! 60 c 0 i E 40 ~ c 8 20 20

I o~~.----.~~::::::::::_____---'"------~- 0 -1----~--~-~------l o 500 1000 1500 2000 2500 3000 3500 4000 4500 0 200 400 600 800 1000 1200 1400 1600 penetration (µm) penetration (µm) Figure I : Interdiffusion profiles

This strong concentration dependence explains the behaviors of the various alloys during their high temperatures nitriding and the diffusion profiles give immediately infonnations about the quality of the future results of the nitriding reactions. When a sharp decrease of the diffusion coefficient is observed in a range of concentrations, the interdiffusion zone is then not wide enough for a good description of the nitriding reaction. If the phase diagram shows phases with limited single phase fields, the related interdiffusion zones must be wide enough for analysis to be possible. For low concentrations in beta stabilizer element, greatest is the size of the atom, lowest seems to be the diffusion of the element. As illustration Table 1 represents interdiffusion coefficient at I 400°C versus atomic radius of the element (for a composition of 5 at.%).

interdiffusion coefficient atomic radius element (5 at.%) 9 2 1 (I 0· cm .s" ) (run)

--- ·-· Nb --- 14.1 0.146

Mo - - 19.6 0.139 v 65.7 0.134 Table I : lnterdiffusion coefficients at I 400°C for low contents of beta stabilizer elements.

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4. NITRIDING OF TITANIUM-X DIFFUSION COUPLES

If Xis a beta stabilizer element (as Mo, Nb or V), nitriding of diffusion couples gives similar results. They.are presented for a nitriding temperanire of l 400°C in figures 2, 4 and 6. Zr is. a "neutrai' element" and gives different results. · For the low concentrations of alloying element nitriding gives the same layers as in the case of the nitriding of pure ti_tanium : an external 8 nitride layer_(fcc structure) with an underlying a solid solution layer (hexagm_lal structure). The difference in the grey level indicates the .0/'8 transition, In- these layers the X concentration is always low and this element is repelled in the core of the sample. This ·effect is .lower for Zr. A low. X concentration (Table 2) is sufficient to slow down the kinetics of the a layer development. For higher beta stabilizer element concentrations, the 8 and a phases expand no more in form of a layer but in form of grains separated by beta metallic matrix zones. Above a typical concentration the nitrogen penetration is strongly slowed down for Mo and Nb, but increased for V. ·

In some samples the innermost part of the external 8-Ti N 1., nitride layer or some precipitates undergo a transformation to E-Ti 2N subnitride (tretragonal structure). The continuity of the titanium profile through this phase shows unambiguously that this phase forms on cooling, because of the too low cooling rate. Its formation ·is a supplementary proofofthe virtually absence of molybdenum, niobium, vanadium or zirconium. · In the core of Ti-Mo or Ti-V samples, the P phase (bee structure) is not transformed during cooling. But in this matrix one can notice very thin precipitates with a thickness lower than 1 µm. X-ray emission images show that they are composed of titanium and nitrogen but the concentrations are not measurable considering their small size. These precipitates probably form during cooling : the P(Ti,X) solid. solution rejects a(Ti) during the monotectold transformatil:m proposed by the Ti-Mo and Ti-V phase diagrams [6].

Temperature of nitriding: - 1400°C Alloying element Mo Nb v Zr Duration of nitriding 6h 3h lhlO 3h

8-(Ti,X)N1., layer mean width (µm) 30 20 20 limits of at.% min 3 8 4 large area* at.% max 10 18 -30 needled area starting composition (at.%) 20 25 cx-(Ti,X)(N) max width (µm) 210 110 86 105 'layer start of width decrease( at.%)· -1 3 4.5 end (at.%) -3 8 9 50 -· ' internal islands higher core cone. (at.%) .-15 20 * correspond with the lower and higher core concentrations where large 8 area is observed.

Table 2. Comparative c~aracteristics for the nitrfding of Mo, Nb,V, Zr. . I . , ' .

4.1. Nitriding of titanium-molybdenum couple (figure 2).

The phases observed and the metallographic features are the same as for nitrided pure titanium if the molybdenum conc.entration is lower than· I at.%. From outside the layers ·observed are 8-TiN 1., and 'a(Ti,N). The molybdenum concentration in these layers is lower ·than 0'.2 at.%. For this reason we have noted 8-TiN 1., rather than 8-(Ti,Mo)N 1., The core of the sample shows a characteristic a microstructure martensitically· transformed from the high temperature p phase during cooling. The nifrogen concentration in this third regiori is lower than the detection limit by EPMA. From I at.% upwards the 8 layer remains unchanged but the a layer thickness is clearly reduced and shows large grains whose boundaries are parallel to the nitrogen diffusion direction. Across the layer, the nitrogen concentration decreases from 22 at.% to 15 at.%. Molybdenum is rejected from these two layers into the P phase

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where its concentration reaches 3 at.% in the first micrometers beneath the a limit while the initial molybdenum concentration was I at.%. In this P phase the nitrogen concentration is lower than the detection limit but nitrogen penetration gives place to internal nitriding by formation of big a(Ti,N) grains.

Figure 2 : Titanium-molybdenum couple nitrided at 1400°C, 6 hours.

For compositions near 2 at.% Mo, the thickness of the 8 nitride layer increases and from 3 at.% Mo to IO at.% Mo the growths of the 8 and a phases are clearly modified. The nitride grains are stretched in the nitrogen diffusion direction and have then a low nitrogen concentration gradient. Backscattered electron images in MEB reveal a thin metallic network (figure 3) decorating the boundaries of these nitride grains. In this metallic network molybdenum concentrations as high as 35 at.% can be detected. This highlights not only the rejection of molybdenum when the nitrogen diffusion forms nitrides but also the short range diffusion of molybdenum which is a consequence of the sharp drop in diffusivity when molybdenum concentration increases. The a(Ti,N) phase grows also as elongated grains parallel to the nitrogen penetration direction. The solubility of molybdenum in these grains is less than 0.5 at.%. When the molybdenum concentration in the initial alloy increases this layer becomes discontinuous and molybdenum rich P phase separates the a grains from one another. The molybdenum concentration in these p grains is high enough (for example 15 at.% whereas the uniform starting value was 2.5 at.%) to stabilize the p phase during cooling. The nitrogen solubility in this phase is very low. The a grains spacing increases as we progress into the sample and finally the initial molybdenum concentration level is found again.

Figure 3 : Molybdenum rich metallic network in the 8 - TiN 1-x layer (uniform concentration in the alloy: 5 at.%)

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For molybdenum concentrations higher than 10 at.% the nitriding kinetics are slightly slowed down. The only nitrided phase is 8-TiN i-x as a continuous layer prolonged by elongated grains into the matrix. The penetration of these elongated grains decreases as the molybdenum concentration increases. The beta phase between these delta grains is enriched in molybdenum with respect to matrix composition in the core of the sample. It must be kept in mind that a good description of the nitriding of Ti-Mo alloys cannot be obtained easily for concentrations higher than 30 at.% when Ti-Mo couples are used. Because of the very sharp decrease of the interdiffusion coefficient when molybdenum concentration increases, the interdiffusion zone between 30 and 100 at.% Mo is not wide enough. For these high molybdenum concentrations experiments must be carried out on Ti,Mo/Ti,Mo or Ti,Mo/Mo interdiffusion samples in order to have wide diffusion zones exhibiting high molybdenum concentrations.

4.2. Nitriding of Ti-Nb diffusion couple (figure 4).

Up to ~ 8 at.%, an observation of two external layers can be done : an external 8-(Ti,Nb) N 1., nitride layer with an underlying a-(Ti,Nb)(N) solid solution layer. In these layers the. niobium concentration is lower than 1.5 at.%. Beneath these external layers, internal nitriding occurs. Large needles and islands are composed of a-(Ti,Nb)(N) solid solution (hexagonal structure). In these islands, the niobium concentration is lower than 5 at.%. For example, the N, Nb and Ti composition profiles measured for an initial niobium composition about 12 at.% are shown in figure 5. Examination of these profiles reveals Nb enrichment of the p matrix, up to 25 at.% beneath the 8-(Ti,Nb) N i-x layer. Niobium is repelled in the residual P phase (bee structure) and the enrichment is higher near the surface than in the core.

A clear distinction between the external 8-(Ti,Nb) N 1., layer and the internal nitriding no longer exists if the initial niobium concentration is greater than 15 at.%. If this concentration is greater than 20 at.% the needles are only composed by 8-(Ti,Nb) N 1., nitride with a niobium composition lower than 2 at.%. The higher the Nb content of the alloy is the smaller are the needles formed during nitriding. EPMA niobium analysis in the p phase reveals a very strong enrichment : 43 at.% for an initial composition of 25 at.% and 63 at.% for an initial composition of 33 at.%. For niobium concentration between 40 at.% and 100 at.%, the shape of the initial diffusion profile shows that our technique cannot have good resolution. These results are in very good agreement with results obtained by Buscaglia [7] and Mayr [8] in their high temperature nitriding of Nb-Ti alloys experiments. Nitriding of an alloy Ti 95 Nb 5 confirms the formation of two layers.

106I at.%Ntl ,\

Figure 4: Titanium-niobium diffusion couple nitrided at 1400°C, 3 hours.

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70 ~ ':; 60 ~ ;:: 50 ~ z 40 x, z- "\ 30 y

0 0 50 100 150 200 250 300 350 400 distance (µm)

Figure 5 : Nitrogen (x), Niobium (0) and Titanium(~) analysis (unifonn concentration in the alloy: 12 at.%)

4.3. Nitriding of titanium-vanadium diffusion couples (figure 6).

Nitriding of the low vanadium concentration area gives an external o-(Ti, V)N i-x nitride layer with an underlying a(Ti,N) solid solution layer. The core remains in the P state and undergoes a martensitic transfonnation during

cooling. EPMA shows that the vanadium concentration in the o-(Ti,V)N 1_xnitride and in the a(Ti,V)(N) layers is low, and lower than 2 at.% V. Vanadium is repelled into the P phase where a concentration increase between J and 2 at.% can be measured right beneath the a layer.

Figure 6: Titanium-Vanadium diffusion couple nitrided at 1400°C, 70 min.

For a concentration about 5 at.% the nitride layer increases and the a-layer decreases. It disappears for a concentration in the order of 10 at.% vanadium. In the 8-layer, small grains are separated by very small zones in which the vanadium concentration is high. This element is rejected from the nitride 8-TiN i-x between its grains but also in the front of the layer : for example, in figure 6 the maximum thickness is observed for a vanadium concentration in the core V-10 at%. The vanadium concentration in the layer it is lower than 2 at.%. but beneath the layer it corresponds to 15 at.% V. The vanadium stabilizes the P phase during cooling and the p phase matrix no longer presents aspects of a martensitic transfonnation after cooling.

In the vanadium concentration area between -25 and -40 at.% V, the 8-TiN 1.xnitride layer thickness is less important but nitrogen penetration gives also internal nitriding. Long islands of 8-TiN l-x nitride fonn with orientation relationships with the P matrix. The 8-TiN 1., nitride layer and islands are directly in contact with the P(Ti,V) matrix.

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If the vanadium concentration is higher than 40 at.%, the external layer becomes a mixed nitride layer (Ti,V)N. It is very thin (a few micrometers) for short nitriding duration. This layer coexists with an internal needle-shaped nitriding that spreads out in a triangular form zone. The thickness of internal nitriding zone increases when the titanium concentration decreases. It is narrowest around 40 at.% and the broadest around 85 at.% vanadium.

This internal nitriding corresponds to 8-(Ti,V)N 1_, needles. They are thus in contact with a p(Ti,V) phase that presents a variable composition. On the titanium rich side needles have a more important thickness with a composition about Ti62.4 - V4 - N33.6 at.%. Between these needles a very rich vanadium p phase (Ti49.5 - V50 - N0.5 at.%) is observed. The needles are very thin at the vicinity of the surface making their analysis very difficult. They seem composed of(Ti,V)N, V2 N 1_y or 8-TiN 1_,: these 3 phases have been identified by EPMA or by X-ray diffraction on bulk homogeneous alloys but remain difficult to localize with certainty. For the high vanadium contents needles become thin near the surface as well as inside the sample. They are dispersed in a vanadium rich matrix. Needles disappear in the solid solution and simultaneously the fine needles starting from the superficial mixed nitride layer transform into real grains : they are V 2 N l-y grains. Gradually the external layer evolves to a vanadium nitride layer and the subnitride V 2 N i-y develops small grains in the layer. When titanium has disappeared in the P solid solution, one can observe by metallographic studies characteristic streaks of V 32 N3 formed during cooling. This phase appears only in absence of titanium. All these results are in good agreement with our results [9] obtained by nitriding of Ti-V alloys.

4.4. Nitriding of titanium-zirconium diffusion couples (figure 7).

The titanium rich side (at.% Zr < 50) of the couple presents an external 8-(Ti,Zr)N 1_, nitride layer whose thickness is uniform (:::020 µm). Beneath this layer, the a-(Ti,Zr)(N) solid solution layer appears whose thickness decreases regularly as a function of the Zr concentration when this concentration is greater than 4 at.%. This layer vanishes when the composition reaches 50 at.%Zr. In opposition to the other alloys, no internal nitriding was observed in the matrix. From these two layers, Zr is repelled into the P phase. Nevertheless, this effect is very smaller than what observed in other alloys. As example, when the Ti/Zr ratio is 3.6 in the core of the P phase, the rejection of Zr forward the nitrogen flux leads to Zr enrichment beneath the a layer (where the Ti/Zr ration is 2.8) and to a Zr depletion in the layers (Ti/Zr=4.5).

Figure 7: Titanium-Zirconium diffusion couple nitrided at 1400°C, 3 hours.

Around the equi-atomic composition of the matrix, the a-layer is so thin that it is not clearly observed. Under the

8-(Ti,Zr)N 1_, nitride layer, an area with no martensitic transformation appears. It seems to be a solid solution of nitrogen in the P-phase of an equi-atomic Ti-Zr alloy. When the zirconium concentration is higher than 50 at.%, an a-(Zr,Ti)(N) appears once again under the external

8-(Zr,Ti)N 1_, layer. Its thickness regularly growths to reach a maximum when composition reaches 100 at.% Zr. From this layers, Ti is partially repelled into the P-phase. When the Ti/Zr ratio decreases in the matrix, it also decreases in the layers, in accordance with the well known mutual solubility of the zirconium and titanium base phases in this concentration range. The rich Ti side alike the rich Zr side seem to have the same reaction with nitrogen : formation of an a solid solution phase beneath a nitride 8 phase where the minor alloy element is weakly repelled into the Pmatrix.

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5. DISCUSSION AND CONCLUSION

These applications of our experimental method for scanning the reaction behavior of alloys with hot gases concern two types of titanium alloying elements. Molybdenum, niobium and vanadium are beta stabilizers.

Niobium and vanadium can react with nitrogen to fonn nitrides, which are less stable than cS-TiN 1.,. On the contrary molybdenum doesn't fonn nitrides in our experimental conditions. This method reveals rapidly that beta stabilizer element (Mo, Nb or V) additions leads to the a-layer thickness decreasing. All the alloying metals are rejected from the phases containing nitrogen and are swept away by the nitrogen flux. Higher is the nitrogen diffusion, higher is the rejection. The short range diffusion of the alloying metal is a consequence of the sharp drop in diffusivity when its concentration increases. Vanadium has the greater diffusivity and allows the better layer growth. Its higher diffusivity for the higher concentrations (V>40at.%) can be an explanation for the original nitriding behavior. Nitriding can only progress by titanium nitriding because of the low nitrogen solubility in the beta phase. But nitrogen diffusivity in the beta phase is better if this phase is a titanium rich phase i.e. ifthe beta stabilizer element is repelled far away from the nitride fonned. When the alloying element is zirconium, which is not a beta stabilizer but a "neutral" element, the results of the reaction with nitrogen are different. No internal nitriding was observed and the nitriding reactions are similar on the both sides of interdiffusion couple (Ti rich side like Zr rich side). Beneath an external layer of cS-(Ti,Zr)N 1., appears a layer of a-(Ti,Zr)(N) whose thickness decreases when the minor alloying element concentration increases. The P phase is only observed after nitriding the equi-atomic TiZr part of the couple. The metallographic and concentration features of the nitrided samples are the results of the composition, the thennodynamic properties and the diffusion characteristics that is the mobility of the species. A judicious choice of the experimental conditions allows an efficient scan of the reaction of alloys with hot gases saving time consuming experiments. Moreover the results can give valuable insights into complicated phase diagrams where various activity levels for the reactive species in the gas phase can be also laid down. This method is not restricted to the study of the four titanium alloys considered in this paper, but can be applied to oxidization, sulfidation, carburation ... of various alloys. Interpretations of the results are nearly simple only when ternary systems composed of two metals and a gas are considered. However this method opens a wide research fields for heterogeneous solid-gas (-fluid) reactions.

6. LITERATURE l. A. GUILLOU, I. THIBON, 0. ANSEL,J.P. BARS,J. DEBUIGNE: "Nitriding of Ti-V and Ti-Mo diffusion couples : examples of a method for scanning the raction behavior of alloys with hot gazes", Proceedings of the symposium on high temperature corrosion and materai/s chemistry, The Electrochemical Society Inc, 1998, 98-9, p.229-240. 2. A. GUILLOU, D. ANSEL, J. DEBUIGNE: "Nitriding of titanium vanadium diffusion couples", Scripta materia/ia, 1998, 38-6, p.981-989. 3. D. ANSEL, I. THIBON, M. BOLIVEAU, J. DEBUIGNE: Interdiffusion in the b.c.c. P-phase of the Ta-Ti alloys, Acta Materia/ia, 1998, 46, p.423. 4. 'r. THIBON, D. ANSEL, M. BOLIVEAU, J. DEBUIGNE: "lnterdiffusion in the p Mo-Ti solid solutions at high temperatures", Zeit Meta//kunde, 1998, 89, p.187. 5. G. LE GALL, D. ANSEL, J.. DEBUIGNE: "lnterdiffusion in the body cubic centered P-phase of Ti-Hf alloys", Acta Meta//urgica, 1987, 35, p.2297. 6. W.G. MOFFAT: Handbook a/Binary Phase Diagrams, General Electric Cy Ed. (1993 update). 7. V. BUSCAGLIA, A. MARTINELLI, R. MUSENICH, W. MA YR, W. LENGAUER: "High temperature . nitriding of Nb-Ti alloys in nitrogen", Journal ofAlloys and Compounds, 1999, 238, p.241-259. 8. W. MAYR, W. LENGAUER, V. BUSCAGLIA, J. BAUER, M. BOHN, M. FIALIN: "Nitriding of Ti/Nb alloys and solid state properties of cS-(Ti,Nb) N", Journal of Alloys and Compounds, 1997, 262-263, p.521- - 528. 9. A.GUILLOU, D. ANSEL, J. DEBUIGNE, M. BOHN: "On the Nitriding behaviour of titanium-vanadium alloys in nitrogen at temperatures between I 150°C and 1400°C", Titanium'95, Science and Technology, 1996, p.2159-2166.

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