Undercooling investigations and heat capacity measurements on PdNiP melts

J O U R N A L OF

Journal of Non-Crystalhne Solids 156-158 (1993) 450-454 North-Holland

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Undercooling investigations and heat capacity measurements on P d - N i - P melts R. W i l l n e c k e r , K. W i t t m a n n a n d G.P. G 6 r l e r Instttut fiir Raumslmulatton, DLR, W-5000 Cologne 90, Germany

The system Pd-NI-P is well-known for its pronounced abdlty to form a metallic glass Studies of the undercoohng behawour of the llqmd alloy were carried out in a DTA facdlty applying the melt-fluxing technique at various coohng rates and overheatmgtemperatures. This technique has been employed to seek the possibd~tyof m SltUinvestigations of the heat capaoty of the undercooled liqmd. The measurements show that crystalhzation can be avoided even for coohng rates as low as 10 K/mln if a sample is exposed several times to overheating temperatures > 1270 K in the B203 flux. This allows for specific heat measurements over a substantial undercoohng range, ff the fluxing method can be adapted for application m calibrated DTA devices Heat capacities, Cp(T), of the alloy m its undercooled llqmd, glassy and crystallized states were measured by standard differential scanning calorimetry and compared with first results using differential thermal analysis techniques on fluxed samples.

1. Introduction Knowledge of the thermodynamic properties of the undercooled liquid is necessary to characterize this metastable state and to describe the solidification behaviour of undercooled melts. In particular, the enthalpy, entropy and Gibbs free energy differences between the metastable liquid state and its stable crystalline phase are of interest. These functions are the basis for a theoretical treatment of all physical processes occurring in the undercooled melt: the nucleation process, phase selection during crystallization, growth velocities, etc. Normally, the thermodynamic functions can be derived from measured heat capacity data. Heat capacities of undercooled liquids, however, are generally not available. This is obviously a consequence arising from the difficulties of handling the undercooled sample in a reproducible manner over sufficiently long timescales. So far, measurements of the heat capacity cover Correspondence to Dr R Willnecker, Institut fur Raumslmulation, Deutsche Forschungsanstalt fur Luft- und Raumfahrt, Postfach 90 60 58, W-5000 Koln 90, Germany. Tel + 49-2203 601 3288. Telefax: + 49-2203 61471.

only a small part of the entire undercooling regime below the melting temperature, Tm, and above the glass temperature, Tg, and have been restricted to low-melting pure metals [1,2] and a few easy glass-forming alloys [3-7], respectively. Various methods appropriate to achieve deep undercooling are applied in laboratory investigations such as containerless processing by levitation techniques, drop tube processing, droplet emulsion technique or the quasi-containerless method of embedding the sample in a matrix of molten glass ('fluxing'). Beside the well-established droplet emulsion method, the last is most suited to carry out the delicate calorimetric measurements, demanding near-equilibrium conditions for extended intervals of temperature and time. However, it has not yet been applied for heat capacity measurements on undercooled glass-forming alloys over the entire metastable temperature range. In the present work, we report undercooling investigations and measurements of Cp on fluxed P d - N i - P samples. The undercooling behaviour of fluxed liquid P d - N i - P ingots was studied in order to establish a suitable thermal treatment that allows maximum undercoolings. Applying the

0022-3093/93/$06.00 © 1993 - Elsevier Science Pubhshers B V All rights reserved

R Wtllnecker et aL / Pd-Nz-P melts

same thermal conditions, first heat capacity measurements were initiated using a differential thermal analysis equipment suited for calorimetric investigation. Further quantitative Cp measurements were performed by differential scanning calorimetry on non-fluxed samples in order to verify the accuracy of obtained data.

2. Experimental details

Primary samples of Pd40Ni40P20 were prepared by alloying the constituents Pd (purity 99.999%) and Ni2P (purity 99.97%) in an arc furnace under clean argon atmosphere. Extended heat treatment under vacuum in a flux of B20 3, followed by a rapid cooling (10-20 K/s) resulted in bulk amorphous samples of cylindrical shape of about 2.2 g in mass. All samples studied in this work were cut from these initial ingots and varied in mass between 30 and 170 mg. The samples were investigated without and with B20 3 embedding. This fluxing material was used to isolate the ingots from the A1203 crucibles and to purify the samples by dissolution of surface impurities. This technique has turned out to be successful in suppressing heterogeneous nucleation of this alloy in previous undercooling experiments [6,8,9]. The composition of the sample after alloying and fluxing and its stability during the undercooling experiments can be monitored by measuring the characteristic temperatures of the melting interval within every thermal cycle; it always ranges from ~ 884 K (eutectic melting) to 986 K (liquidus temperature) according to the phase diagram of this alloy [12]. The mass loss of the sample after an extended sequence of temperature cycles in B20 3 serves as an additional proof; the difference to the initial mass did not exceed a value of 10 I~g. The undercooling experiments were performed in a differential thermal analysis (DTA) equipment (Netzsch DTA 404S) with an operational temperature range up to 1600°C, in a heat-flow DSC facility (Netzsch DSC 404) with an upper limit at 1500°C and with a power-compensating differential scanning calorimeter (Perkin-Elmer DSC-2C), restricted to temperatures below 1000 K. During the experiments, the facili-

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ties were constantly purged with He and Ar gas, respectively.

3. Results

The undercooling behaviour of the PdNiP alloy was investigated by exposing the liquid samples embedded in B20 3 to different overheating temperatures prior to the subsequent cooling cycle, varying the holding time and the maximum temperature. From previous undercooling studies on glass-forming systems, it is known that the nucleation rate can be essentially influenced by the overheating temperature and time, thus favouring amorphous solidification also at low cooling rates [10,11]. In the present work on PdNiP alloys, the overheating temperature was step-wise increased starting from T = 1036 K to T = 1450 K. Choosing a constant overheating time of 10 min and a constant cooling rate of 10 K/min, we found in general decreasing solidification temperatures for the onset of solidification (fig. 1). For moderate overheating, the near-equilibrium solidification is observed with primary crystallization of NisP 2, continuous solidification within the liquidus range and a prominent peak due to the eutectic solidification at or below the eutectic temperature. When increasing the overheating temperature, both peaks are shifted towards lower temperatures, finally coinciding into one solidification event on the DTA trace at a low temperature level. Repeated cycles of the fluxed melt up to temperatures of 1270 K or higher result in a further reduction of the nucleation. During the subsequent cooling phase, the liquid sample can freeze into the amorphous state, showing only weak indications of a hindered crystallization between 710 and 650 K. From the residual crystallization energy, a crystalline fraction of 1 or 2% can be concluded. In many cases, the cooling curve is without any crystallization exotherm. Figure 1 shows the corresponding curve, thus demonstrating that, even for cooling rates of 10 K/min, crystallization can completely be avoided for this alloy. Comparing this result with previous observations on P d - N i - P [9], the critical cooling rate necessary to prevent crystal-

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lization can be reduced by nearly one order of magnitude merely by application of an appropriate thermal treatment of the liquid sample. It became obvious from our experiments that primarily the purification of the sample is responsible for the disappearing of any crystallization, since - after a few thermal cycles - freezing into the amorphous state occurred even following a moderate overheating temperature as low as 1160 K. Differential thermal analysis was applied to prove the glassy state of the ingots prepared at the slowest cooling rate of 10 K/min. Figure 2 shows a typical DTA trace (heating rate 20 K/min). The curve includes all characteristic ki-

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netics: the transformation from glass to undercooled liquid (glass transition temperature, Tg--583 K), the transition into the crystalline state (crystallization temperature, T c = 668 K) and finally the melting. Melting of the sample occurs between 884 and 975 K, in accordance with the equilibrium phase diagram [12]. The crystallization energy of AH c = 93.2 J / g resulting from our measurements is in good agreement with previous findings on rapidly solidified samples [6]. From this, it can be concluded that the sample consists at least of 95% amorphous phase. At a temperature of about 785 K, the beginning of an additional weak endothermic transformation is observed with a peak temperature of 807 K, not mentioned in previous investigations. With respect to the value of A H c, the energy AHtr = 3.5 -I- 0.2 J / g of this peak is in the order of a crystalline phase transformation. The peak occurs not only following the crystallization of glassy samples, but also during the heating phase of samples solidified from undercoolings down to at least 770 K. The origin of this weak endotherm has not been clarified unambiguously and will be the subject of further investigations. Heat capacities of non-fluxed samples were measured in the DSC equipment. The heat capacity of an amorphous sample from the glassy state to the undercooled liquid above Tg is shown in fig. 3, together with that of a crystallized sample. The amorphous sample was several times cycled around the glass transition with decreasing cooling rate down to 5 K/min. During this treatment, the temperature did not exceed 625 K, to prevent any spurious crystallization. The crystalline specimen was annealed at 850 K before measurement to approximate the equilibrium state. Within the limits of uncertainty, the obtained values are in quantitative agreement with previous results [6]. At temperatures T < 550 K, the Cp curve of the glass approaches that of the crystal and intersects this curve at T = 480 K. This is caused by the magnetic contribution to the heat capacity of the crystalline state manifesting itself by the magnetic transition a t Tmag = 513 K. This contribution has not yet been reported in calorimetric measurements on this system, but is well known from detailed magnetic investigations

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Pd4oNl4oP20 Curves a and b are obtained from standard differentialscannmg calorimetric(DSC, Perkin-Elmer). Curve c is obtamed applying dffferentmlthermal analysis techmques on a fluxed sample Note that samples used for curves b and c were cycled around Tg with the same coohng rate of 5 K/mm before measurement

453

described above offers the possibility to determine in situ the specific heat of the fluxed samples in this metastable liquid state. The achieved low critical cooling rate is compatible with the rates of temperature change used in heat-flow DSC equipment for calorimetric investigations. Thus, the entire temperature range of the undercooled melt down to Tg should become accessible for Cp measurements. However, as a necessary precondition the heat capacity of the fluxing material surrounding the sample must be compensated or taken into account by an appropriate method. With this restriction, the fluxing technique can be adapted for use in a conventional heat-flow DSC. Three different scans have to be compared: (1) a crucible filled with B203, (2) a sapphire, embedded in B203, as a standard for calibration, (3) the metallic sample enclosed in B20 3. As reference, a crucible filled with equal mass of B 2 0 3 is recommended. Testing this procedure first successful measurements have been performed on an amorphous sample in the vicinity of the glass transition. The data obtained are shown in fig. 3 as a dashed line (curve c), fitting quantitatively very well with the data obtained from DSC.

5. Conclusions

[12]. With mcreasing temperatures, the curves diverge markedly, indicating the glass transition at Tg = 583 K. The specific heat of the undercooled liquid can be measured over the subsequent temperature range up to approximately 635 K, before it is interfered with by the weak onset of crystallization. The difference in heat capacity between the undercooled liquid and the crystalline solid is determined to A C I - ' ( T g ) = 0.225 J / g K, and the difference in Cp between the undercooled liquid and the amorphous state is obtained as = 0.153 J / g K.

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4. Discussion

The accessibility of the large undercooling range achieved by the experimental procedure

We have investigated the undercooling behaviour of Pd40Ni40P20 alloys by using the melt fluxing technique in a D T A facility. The experiments were aimed to find out an appropriate thermal treatment of the sample to reach undercooling levels down to the glass temperature. This gives the prospect of performing in situ measurements of the heat capacity of the metastable liquid. It could be shown that heating the sample in B20 3 embedding to a certain temperature (T ~ 1270 K) above the melting point for a sufficient long time gives rise to the prevention of crystallization completely even for cooling rates as low as 10 K / m i n . This allows for specific heat measurements in a conventional heat-flow DSC equipment if the heat capacity of the fluxing substance can be taken into account in a suitable manner. We have demonstrated the feasibility of

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R, Wdlnecker et al. / P d - N t - P melts

this method by comparing the results for Cp of fluxed samples obtained in the differential thermal analyzer with DSC measurements on samples with the flux removed. Further experiments are expected to cover the entire undercooling regime of a metallic system for the first time. The authors want to thank Mrs Heike Miihlmeyer for sample preparation and Mr Gerhard Wilde for performing the DTA measurements.

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[3] H,S. Chen and D. Turnbull, J Appl Phys 38 (1967) 3646 [4] H.S. Chen and D. Turnbull, J. Chem Phys. 48 (1968) 2560. [5] P.V Evans, A. Garcla-Esconal, P E. Donovan and A.L Greer, Mater. Res Soc. Symp. Proc 57 (1987) 239 [6] H W Km and D. Turnbull, J Non-Cryst. Solids 94 (1987) 62. [7] M C Lee, H.J. Fecht, J.L. Allen~ J H. Perepezko, K Ohsaka and W.L. Johnson, Mater. SCL Eng. 97 (1988) 301. [8] A J Drehman, A.L. Greer and D. Turnbull, Appl. Phys Lett 41 (1982) 716 [9] H.W. Kin, A L Greer and D. Turnbull, Appl Phys. Lett. 45 (1984) 615. [10] F Gdlessen, D M. Herlach and B Feuerbacher, Z. Phys. Chem 156 (1988) 129. [11] Z. Altouman and J O Strom-Olsen, m' Proc. 6th Int Conf on Rapidly Quenched Metals, Wurzburg, 1984 (North-Holland, Amsterdam, 1985) p. 447, [12] E. Wachtel, H. Haggag, J Godecke and B Predel, Z Metallkd. 76 (1985) 120