Effect of the addition of palladium on grain growth kinetics of pure titanium

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AHD C£)MPG~L~3 ELSEVIER

Journal of Alloys and Compounds 260 (1997) 147-152

Effect of the addition of palladium on grain growth kinetics of pure titanium F.J. Gil a'*, J.A. P i c a s b, J.M.

Manero ~,

A. F o r n b, J.A. P l a n e l l a

~Universidad Polit(cnica de Cataluna, Dept. Ciencia de los Materiales e Ingenieria MetaMrgica, E.T.S. Enginyers Industrials de Barcelona, Avda Diagonal 647, 08028 Barcelona, Spain bE.U. Politbcnica de Vilanova i la Geltrd, Universitat Poiit~cnka de Catalunya, Barcelona, Spain

Received i7 February 1997; received in revised form 14 March 1997

Abstract In the present work, experimental research on grain growth kinetics at different temperatures and times of heat ~eahment for titaniam and Ti-0.2Pd in the cx and [3 phases has been carried out. The grain size parameters were obtained by the image analysis technique. The aim of the present study was to quantify the progress of grain growth and the corresponding activation energy for ~xand [3 titanium, and to determine the effect of Palladium. By means of EDS microanalysis, it has been observed that the effect of a higher concentration of palladium on the grain growth is to decrease the growth exponent for the Ti-0.2Pd alloy. The activation energies differ a lot between the two phases oLand 13for Ti and Ti-0.2Pd. This difference might be due to the different crystal structures, © 1997 Elsevier Science S.A. Keywords: Grain growth; Titanium; Heat treatment

1. Introduction The technological importance of grain growth stems from its effect on various properties, and in particular on the mechanical behaviour. In materials for structural applications at lower temperatures, a small grain size is normally required to optimise the strength and toughness. A good understanding of grain growth is therefore a prerequisite for control of the microstructures and properties of metals and ceramics during solid state processing

[ll. The driving force for this process results from the decrease in free energy which is due to the reduction in total grain boundary area [2]. Grain growth takes place by diffilsion when the temperature is high enough and the time of heat treatment is long enough. This means that the number of grains/unit v,~lume decreases, the size of the grains increases, and both the grain boundary area/unit volume, and the stored energy/unit volume, decrease. Consequently a state of higher thermodynamic stability is reached. This process of grain growth can occur by two mechanisms: normal grain growth and abnormal grain growth or secondary recrystallization. During normal grain growth, the microstructure changes in a rather uniform *Corresponding attthor. 0925-83881971517.00 © 1997 Elsevier Science S.A. All rights reserved. Pll S0925-8388(97)00135-7

way. There is a relatively narrow range of grain sizes and shapes, and the form of the grain size distribution is usually independent of time. During abnormal grain growth, a few grains in the microstructure grow and consume the matrix of smaller grains and a bimodal grain size distributnon develops [1-4].

2. Experimentan procedure This study was carried out with commercially-pine titanium (sample A) and a Ti-0.2Pd alloy (sample B) which has been kindly donated by Technalloy S.A. in the form of cylindrical rods of 6 mm in diameter and 200 mm in length. Disks 6 mm high were cut out and 38 specimens were prepared from both samples. The chemical compositions of these alloys are shown in Table 1. The microstructure of the material, as received, consisted of equiaxed grains of the cx phase for samples A and B. The microstracture of the latter is shown in Fig. 1. Table l Chenfical compositions in weight percent Ti Ti-0.2Pd

PaIFadiura Ni~ogen Carbon Hydrogen Iron

Oxygen

0.18

0.10 0.15

0.005 0.007

0.01 0.02

0.96 0.02

0.03 0.05

F.J. Gil et al. I Journal of Alloys and Compounds 260 (1997) 147-152

148

Fig. 1. Microstructure corresponding to a transversal cut of the rod of Ti-0.2Pd (material as received).

Two samples were used as reference, while the others were subjected to different heat treatments at temperatures of 700, 750, 800, 900, 1000 and I I00°C by applying heating times of either 3, 5, 10, 15, 30, 60 or 120 min at each temperature. The u and 13 phase grain growth kinetics were studied. The u~13 transitions occurs above 882 °C for the pure titanium (sample A) and between 790 to 913 °C for the Ti-0.2Pd alloy (sample B). A set of specimens was placed into the furnace at a fixed temperature for each experiment and then removed from the furnace after an appropriate time of heat treatment and quenched into water at 20 °C. This procedure gave the same cooling rate for all the samples. Subsequently, the specimens were metallographically polished and etched with a solution containing both HF and HNO 3. Metallographical observation was carried out by optical microscopy. The grain size parameter was obtained by the Image Analysis Technique with the Omnimet 3 program. The image was enhanced in contrast and pseudocolor in order to facilitate its interpretation. After this optimization process, grain growth parameters were ideatified and quantified. Finally, the image was coded by a computer and the data were statistically analysed.

3. Results and discussion The measured grain size as a function of the time and temperature of heat treatment is plotted in Fig. 2a Fig. 2b for the ot and 13 phases of sample A, respectively. In Fig.

3a Fig. 3b results are shown for the ot and 13 phases of sample B, respectively. As expected, on raising Lh.e temperature the kinetic growth rate is faster and on extending the period of time at the same temperature, grain growth is greater. From Figs. 2 and 3 it is possible to notice that the grain growth takes place at a very fast rate in the first 10 to 15 rain of the heat treatment and for such times onwards the growth rate decreases. This rate decrease can be explained by taking into account the grain size increase which produces a decrease in the grain boundary area per unit volume ratio. This might mean that the grain boundary inteffacial energy per unit volume decreases and therefore, the driving force for grain growth is lower. During the grain growth process, the equiaxed grains are clearly observable. Since the grain size in the longitudinal and in the transverse direction appeared to be the same further grain size measurements were performed only by examination of the cross-section of the extruded bars. This means that the texture effect is probably not significant because the microstructure of extruded samples showed complete recrystallization. Fig. 4 shows the pure titanium sample heat treated at 1100 °C for 120 rain. Considerable grain growth can be observed. The microstructure corresponds to the ~t' martensite with an acicular morphology. Kinetic grain growth follows Hillert's distribution since the maximum radius is 1.8 times larger than the average radius value. This means that uniform growth occurs in the whole sample and size distribution obeys an asymptotic law typical for a state of equilibrium [5]. When the grain size data were plotted as a function of log D versus log t a

F.J. Gil et al. I Journal of Alloys and Compounds 260 (1997) 147-152

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Fig. 2. (a) Ratio of diameters in relation to temperatures and times of heat treatment for pure Ti in a-phase and a---->[3transition. (b) Ratio of diameters in relation to temperatures and times of heat treatment for pure Ti in [3 phase.

straight line resulted. This means that the generM expression for grain growth can be given by the following equation [3,4]. (1)

where D is the mean grain size, D O is the initial grain size, K is a tew:perature dependent constant, t is the time and n is the time exponent obtained from the slope of the log D vs log t plot. Moreover, if the assumption is made that atomic diffusion across the grain boundary is a simple activated process, and when n is independent of the temperature, h follows that K can be wfiuen as [6,7]: K = Ko e x p ( ~ )

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where E~ is the activation energy, T is the temperature in Kelvin, K o is the we-exponential rate constant and R is the gas constant. This would explain why a plot of L n ( D - D o) vs l i T produces a straight line. In that case the slope

Fig. 3. (a) Ratio of diameters in relation to temperatures and times of heat treatment for Ti-O.2Pd in a-phase and a ~ B transitions. (b) Ratio of diameters in relation to temperatures and thnes of heat treatment for Ti-O.2Pd in [3 phase.

sho,Jld be - E a IR. In Fig. 5 the linear relationship between l n ( D - D o) and 1/T for grain growth of ~ phase of the Ti and Ti- 0 .2 P d can be observed. The kinetic behaviour is in agreement with these assumptions, since after taldng logarithms, linear equations are obtained with good correlation coefficients. The growth exponent n for each hetat treatment temperature is shown in Table 2. From these data, it can be derived that the growth exponent increases with temperature. However, these values are higher than those found for other metals and alloys [8-10]. in the ideal case the growth exponent would be 0.5 but generally n is observed to be less th,'m 0.5 due to the role played by different grain growth parameters such as impurity-drag, free surface effect, texture, dislocation substructure and heterogeneities [11,12]. In our case, the titanium studied presents high purity and very low dislocation density because it is a completely recrystallized metal. Probably this is the reason that the growth exponent is around 0.5. However, the final average grain size for the specimens treated at 900 °C showed no increase of area for the a phase, due to the

150

F.J. Gtl et al. I Journal of Alloys and Compounds 260 (1997) 147-152

Fig. 4. Martensitic microstructure of pure tilaniu, n heat treated at 1100°C for 120 rain.

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From results of Table 2 and Table 3 it can be derived ~hat, for file Ti-0.2Pd alloy the grain growth kinetics is ,;omewhat different than for the pure titanium. It is well known that a grain boundary often has a different composition than the surrounding grains. The difference in Pd concentration between a high angle grain boundary and the matrix has been determined by means of EDS microanalysis. It can be seen in Fig. 7 that the Pd content in the grain boundary is about 1.1% and it decreases rapidly to 0.1% in the matrix. The effect on the grain growth of a higher concentration in the grain boundaries will be a 'solute drag effect'. This mechanism assumes that the movement of the boundary is determined by the diffusion of solute elements behind the boundary. A grain boundary moves with a velocity (v) in response to the net pressure (P = ~vPi) on thc boundary. It is generally assumed that the velocity is directly proportional to the pressure, the constant of proportionality being the graia boundary mobility (M) and thus,

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proximity nf the temperature to transition point. Apparently the thermal energy was used mainly for the nucleation of the 13 phase in ot grain boundaries. Consequently, the diffusion process that produces the grain growth is inTable 2 K. n values for different treatment temperatures for Titanium c.p.

v =MP

T

log K

n

700 750 800 900 1000 i 100

0.03 0.98 0.98 0.06 3.41 4.2

0.30 0.35 0.50 0.05 0.60 0.50

(3)

This type of relationship is predicted by reaction rate theory if the mobility is independent of the driving force and if P << kT, and it should be independent of details of the of boundary migration mechanism. If the boundary mobility is controlled by the s~.~ime atoms, the retarding pressure due to solute drag (P~o~) may be a function of the boundary velocity. Althougta a

F.J. Gil et al. i Journal of Alloys and Compounds 260 (1997) 147-152

,I ~~a

Fig. 6. Initial phase transformation mainly on the grain lmundaries for a sample of pure titanium.

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complex relationship between driving pressure (Pd,) and velocity may be found, equation [3,4] may still be obeyed if P is replaced by (Pd,-P~o~) [13-15]. The Ti-O.2Pd alloy, when treated at 900 °C showed no increase of the area of the ot phase, due to the proximity of Table 3 K, n values for different treatment iemperatures for Ti-0.2Pd T

log K

n

700 750 800 900 i000 1100

0.05 0.20 |. 10 1.20 2.35 2.65

0.15 0.16 0.05 (J.07 0.30 0.25

the a--->13 transition temperatures. As with the pure titanium, the thermal energy was probably used mainly for the .III.U. L. '^-+:"" ~" "" h . . . . a ~ , ~ n,, IIK~alI.IUI 1 of the 13 "i" "~ " +~° ~taa Fig. 8 one can observe the initial phase transformation in the grain boundaries. The activation energies differ a lot between the two phases cx and 13 in each alloy Table 4. The difference might be due to the difference in crystal structure. The ~ phase has a hexagonal structure and the 13 phase has a body centered cubic structure. The |atter structure is more open than the hexagonal one, prone to diffusion and iherefore expected to result in a higher growth rate. The effect of solute additions on the recrystallization of titanium has been studied by Reinbach and Nowikow [ ~6]. Of the solutes added (iron, aluminium, tantalum, tin and chroplium) chromium was found to have the greatest effect in reducing the recrystallization rate. This fact explains the low values of the activation energy in commercially pure titanium. Reinbach and Nowikow found for the [3 phase of the Ti-6A1-4V alloy an activation energy of 95 KJ/mol [17], for Ti-6A1--5Zr-0.5Mo-0.25Si 203 KJ/mol [18], and for the T1-5.6AI-3.5Zr- ~Nb-0.z5Mo-0..~ol 211 KJ/ tool [ 18]. The activation energy for diffusion has a val~e higher than the activation energy for self-diffusion in grain Table 4 Activation energies in KJ/mol

tx

Ti c.p.

Ti-(L2Pd

1OO 21

133 56

152

F.J. Gil et al. I Journal of Alloys and Compounds 260 (1997) 147-152

Fig. 8. Initial phase transformation on the grain boundaries for Ti-0.2Pd.

boundaries. If the atoms of the alloying element would have had the same size as the matrix atoms, the activation energy would probably have been suitable for grain boundary diffusion. The differences observed between the Semiatin's model [19] for beta grain grewth in titanium alloys and our resu!ts for grain growth kinetics is due to the good accaracy of the Analysis Image Technique and the computational data analysis. Others studies [13,20] reported deviations from the theoretical model using a line intercept technique for determining the grain growth kinetic.

Acknowledgements The authors are grateful to Technalloy S.A. for kindly donating the material.

References [1] F.I. Humphreys, M. Hatherly, Re,crystallization and Related Annealing Phenomena, Elsevier, Oxford, 1995, p. 2S2.

[2] H.V. Atkinson, Acta Metall. 36(3) (1988) 469-491. [3] D. Weaire, N. Rivier, Contempory Phys. 25 (1984) 59. [41 F. Haessner (Ed.), Recrystallisation of Metallic Materials, Dr. Riederer Verlag, GmbH, Stuttgart, 1978. [51 M. Hillert, Acta Metall. 13 (1965) 227. [61 I.M. Lifschitz, V?¢.Slyozov, J. Phys. Chem. Solids 19 (1961) 35-50. [7] B.A. Beck, J. Appl. Phys. 19 (1948) 507. [81 A.G. Guy, LJ. Hren, Elements of Physical Metallurgy, 134, Adison Wesley, Reading, MA, 1980. [9] R.L. Fullman, American Society for Metals, Seminar Metals Interfaces 179 (1952) 54. [10] J.C. Fischer, R.L. Fullman, Progress in Metal Physics, Chalmers ed. 97, 1952. [ll] G.T. Higgins, Metal Sci. 8 (1974) 143. [12] R. Elst, J. Van Humbeeck, L. Delay, Z. Metallhde. 76 (1985) 705. [13] H. Frediksson, Mat. Sci and Technol. 6 (1990) 811-817. [14] B. Ralph, Mat. Sci. and Technol. 6 (1990) !!39. [15] F.J. Gil, J.A. Planell, Scripta Met. et Mat. 25 (1991) 2843-2848. [16l R. Reinbach, A. Nowikow, Z. Metallkde. 47 (1956) !196. [17] F.J. Gil, E Tarin, J.A. Planel!, m: F.H. Froes, Y. Captav. (Eds.), Titanium 92 Science and Ter.l~aoiogy, TM ~, 1993 pp. 771-782. 1181 S.P. Fox, Titanium 92. Science and Techn,~logy, in: F H. Froes, 1. Captan (Eds.), TMS, pp. "69-776 (1993). [191 S.L. Semiatin, J.C. Sov~r, I.M. Sukonnik, Scriptii Met. et Mat. 30(7) (1994) 95 ! -955. [20] E.P. Abrahamson, Trans. Met. Soc. AIME 221 (!961) 1193.