- Email: [email protected]
Magnetothermal study of nanocrystalline particle formation in amorphous electroless Ni–P and Ni–Me–P alloys
Electrochimica Acta 47 (2001) 359– 369 www.elsevier.com/locate/electacta
Magnetothermal study of nanocrystalline particle formation in amorphous electroless Ni–P and Ni–Me–P alloys D. Tachev, J. Georgieva, S. Armyanov * Institute of Physical Chemistry, Bulgarian Academy of Sciences, Sofia 1113, Bulgaria Received 1 October 2000; received in revised form 13 March 2001
Abstract The microstructure of amorphous Ni–P and Ni–Me– P materials and especially its change during the heat treatment is of great importance for their magnetic, mechanical and corrosion behavior. A new magnetic phase analysis method (magnetothermal) is presented that reveals the precipitation of nanoparticles with strong magnetic properties during phase transformation upon heat treatment. It is applied to electroless Ni–P, Ni–Cu–P and Ni– Sn – P amorphous alloys. The results acquired by this method are compared with data obtained by differential scanning calorimetry, as well as by microhardness measurements using identical heat treatment in all three cases. Due to the high sensitivity of the magnetothermal method a more detailed picture of the precipitation processes in Ni–P alloys is obtained and the new information is discussed. Magnetothermal measurements reveal several stages of precipitation of a phase with strong magnetic properties. This phase is Ni in the Ni– P alloy, and Ni(Me) solid solution in the Ni–Me–P alloys. Though Sn has a stronger effect on the Ni magnetization, Cu is more effective in preventing the appearance of high magnetization in a thermally treated Ni–Cu–P alloy. This is due to Cu incorporation in Ni particles in a quantity above four times larger than Sn. © 2001 Elsevier Science Ltd. All rights reserved. Keywords: Ni– P, Ni– Cu–P and Ni–Sn–P alloys; Crystallization; Magnetic properties; Microhardness; Differential scanning calorimetry
1. Introduction Electrolessly deposited (ED) Ni – P alloys are unique engineering materials. With the now growing interest in nanomaterials the ED Ni – P alloys remain in focus. Low phosphorus alloys are nanocrystalline as plated [1], whereas high phosphorus alloys are brought to nanocrystalline state by crystallizing an amorphous alloy [2]. Ni –P coatings have a bright and shiny appearance, outstanding mechanical properties and they adhere strongly to the substrate. They possess excellent wear and corrosion resistance. The ED Ni – P alloys with different composition find applications in many branches of industry from food and textile to automotive and aircraft [3]. For example, ED Ni – P utilization in thin film disks is well known [4] and its possible use in flip-chip packaging is currently under investigation [5]. * Corresponding author. E-mail address: [email protected] (S. Armyanov).
The crystallization of Ni –P amorphous alloys has long been studied by electron diffraction, X-ray diffraction (XRD), transmission electron microscopy and differential scanning calorimetry (DSC) combined with other methods including microhardness and magnetic measurements [6–17]. The overall crystallization process can be summarized as follows [9–11,18 –20]: for hypoeutectic alloys (in the range 15 –19 at.% P) a reaction of the primary crystallization of Ni is observed at low temperature and a eutectic transformation, with Ni and Ni3P as final products, at high temperature. Only a eutectic reaction occurs in alloys with eutectic composition (19 –20 at.% P). In hypereutectic alloys, Ni, Ni3P and Ni5P2 crystalline phases form at low temperature. At higher temperature the metastable Ni5P2 phase transforms into Ni3P. All these reactions are present as distinct peaks in the DSC thermograms. The peak position dependence on phosphorus content is summarized in Refs. [9–11]. Ternary alloys, Ni –Me –P, based on the Ni –P binary system were introduced because of the necessity of increasing the thermal stability of the paramagnetic
0013-4686/01/$ - see front matter © 2001 Elsevier Science Ltd. All rights reserved. PII: S 0 0 1 3 - 4 6 8 6 ( 0 1 ) 0 0 5 8 7 - 4
360
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
properties of amorphous Ni– P alloys used as an underlayer in thin film memory disks [19,21]. Nontransition elements have been added to the Ni– P alloys as a third component namely Cu [10,18– 21], Sn [18,22] and Sb [18]. It was shown that Ni– Cu alloy’s particles form in the Ni – Cu –P alloys and this decreases their magnetization after heat treatment [18,19,21]. Both Cu and Sn increase the crystallization temperature and the crystallization activation energy and diminish the alloy magnetization after heat treatment [18]. The highest thermal stability of the amorphous state was achieved by the addition of Sb, but the magnetization of the annealed samples was also high [18]. Magnetic measurements are generally used when a certain magnetic property is the target or for investigating the origin of certain magnetic behavior. Despite the fact that magnetic phase analysis has been known for a long time [23,24] magnetic measurements are rarely used for the characterization of sample structure and phase transitions. This is due to the difficulty in obtaining quantitative information. Magnetic phase analysis is based on the fact that the magnetic moment of a sample consisting of several phases is a sum of their moments and each phase magnetization has characteristic temperature dependence. The aim is to find the mass or volume fraction of each phase. The method is applicable to phases with strong magnetic properties. Masui et al. [15] performed magnetization measurements during the heating of amorphous Ni– P samples. Since the sample is initially paramagnetic and Ni precipitates during crystallization, the stages of crystallization were traced. Similar experiments were made by measuring magnetic susceptibility instead of magnetization [25–29]. Lu¨ ck et al. [27,28] applied the Johnson– Mehl –Kolmogorov – Avrami analysis to the susceptibility curves. However, their assumption [28] that the susceptibility is proportional to the volume of the crystalline phases is rather presumptuous, since the magnetic response of nanometer size Ni solid solution particles is essentially nonlinear. A better approach to the investigation of small particles of strong magnetic phases in a para- or diamagnetic matrix is superparamagnetic (SPM) granulometry [30,31]. It is capable of finding magnetic particles’ size distribution, as well as their volume fraction. The basic property behind magnetization and susceptibility is the magnetic moment. It is a function of both the magnetic phase volume and composition. This is a significant difficulty especially for studies of phase transitions involving simultaneous change of the volume and composition of a phase with strong magnetic properties. Determination of the chemical composition by an independent method is required for phase identification. Magnetic measurements, however, have a number of advantages. First, we should mention the sensitivity.
Magnetization measurement with a standard vibrating sample magnetometer (VSM) could be up to hundred times more sensitive to the volume change of a Ni phase in the Ni–P sample than DSC. Thus the magnetization measurements are a better choice for early detection of a transformation involving phases with strong magnetic properties [28]. Second, the magnetic response is selective. It is very large for phases with strong magnetic properties compared to phases with weak magnetic properties. Thus, only the formation and disappearance of phases with strong magnetic properties will be discerned contrary to DSC, that will detect any transformation connected with substantial heat release. Third, the magnetic response is practically immediate. This is particularly important for the investigation of kinetics. In order to receive more comprehensive information for the crystallization of amorphous Ni–P besides the abovementioned combination of several techniques, the necessity of introducing new methods is obvious. Recently, the results of the use of atom-probe field ion microscopy in combination with DSC and XRD have been summarized [1]. In our previous work [32], we described a new magnetothermal method appropriate for approximate determination of the total volume of the SPM and ferromagnetic (FM) particles precipitating in amorphous Ni–P alloys. In this paper, we will use the magnetothermal method to reveal more details about the formation of nanocrystalline particles in the course of the crystallization process of the Ni–P, Ni – Sn–P and Ni –Cu –P alloys. The results will be compared with DSC and microhardness measurements applying identical heat treatment.
2. Experimental The samples used in this investigation were amorphous ED foils of the Ni1 − x Px alloy (after dissolution of the substrate), as well as melt-quenched (MQ) ribbons and ED foils of the Ni1 − x − y Mey Px alloy, with 16B xB 20 and 0.2ByB 3.0. Electroless (autocatalytical) deposition was performed in Ni sulfate solution, using hypophosphite as the reducing agent. The chemical composition of the plated alloys was determined using energy dispersive spectrometer on a scanning electron microscope JEOL 733. The changes of magnetization at field strengths of 8– 12 kOe were measured while heating the sample from 420 to 720 K at a constant rate of 3 K/min (isochronal heating). Samples were subsequently heated faster to 800 K and held there for about half an hour to completely crystallize them. The magnetization was also recorded during subsequent cooling at 12 K/min. A second cycle of heating and cooling was performed to assure that no further crystallization and accompanying
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
change in magnetization were present. This confirmed that the samples crystallized completely for half an hour at 800 K. All measurements were performed on a VSM, Riken Denshi BHV-55, under vacuum better than 10 − 4 Torr. The DSC measurements were performed in a Perkin–Elmer DSC-7 using N2 as a purge gas. The mass of each sample was 3– 12 mg. The scan rates were identical to those of the magnetothermal measurements. Samples for microhardness measurement were heated with a constant heating rate of 3 K/min in a programmable oven. They were put in the oven at 390 K, pulled out at different temperatures and quenched in cold water. The microhardness was tested by the Vickers method. The load of 245.2 mN was applied for 20 s. To ensure that the thickness of our samples was enough for the correct hardness measurement with this load the softest samples were measured also at a load of 147.1 mN. The same microhardness value was obtained.
3. Results
3.1. Magnetic beha6ior of Ni particles depending on their size The Ni – P alloys represent most accessible case for magnetic phase analysis of precipitation of particles of strong magnetic phase in a matrix with weak magnetic phase properties. The magnetic behavior of pure Ni particles strongly depends on their size and systematization of this behavior based on the data from the literature is presented here. Four groups could be considered. Starting from the smallest size they are giant moment paramagnetic (GMP) clusters, SPM, single domain (SD) and FM particles (Fig. 1). Coupling between atomic magnetic moments, that is characteristic for strong magnetic phase, exists in all four types of particles. The GMP clusters consist of a small number of atoms. The magnetic moment of the cluster is a sum of the moments of the atoms. According to quantum mechanics the magnitude and the projection of this moment can take only discrete values. Ni particles with linear dimensions less than 4 A, behave like GMP
Fig. 1. Size dependence of the magnetic behavior of isolated (magnetically non-interacting) Ni particles at 300 K.
361
clusters [33,34], while particles of 7–8 A, are described as SPM [35]. Due to exchange interaction the SPM particles have significant magnetic moment, but their size is still too small to keep it constant. Thus, because of the thermal agitation their magnetic moment changes its direction and magnitude, so that no remanent magnetization and coercive force can be observed. The magnetic moments of these particles behave like moments of atoms of a paramagnetic material. The upper limit for SPM behavior of Ni particles at 300 K is about 188 A, . With increasing particle size the thermal agitation can no longer disorder the atom moments in the particle, so a hysteresis loop is formed. However, the size of these particles is comparable to the magnetic domain wall thickness and domains cannot be formed. Thus, the particles are called SD particles. The SD particle is always magnetized to almost complete saturation [36]. Two modes of magnetization reversal in this type of particles are possible. A coherent rotation of the magnetization is observed, when all atomic moments in the particle rotate uniformly, and incoherent rotation occurs, when atomic moments do not remain parallel during magnetization reversal. With further increase of particle size its volume becomes large enough to contain more than one magnetic domain. The particle becomes multidomain like most bulk FM bodies. The rigorous mathematical consideration shows that a range of dimensions can mark the transition from SD to MD behavior, rather than a single value [37]. Despite this, the value of 382 A, is given for Ni particles. From equations presented in Ref. [38] the range of 312– 1260 A, can be estimated for Ni at 300 K. It must be noted that all boundary dimensions are temperature dependent. Fig. 1 gives the size dependence of the magnetic behavior of Ni particles at 300 K. Mostly SPM particles are encountered when investigating the crystallization of the Ni– P amorphous alloys.
3.2. Magnetothermal, DSC and microhardness measurements of Ni–P alloy with different compositions The first step in the magnetothermal method is to measure the magnetization at constant magnetizing field strength during isochronal heating. Then the change of magnetization is a result of the combined action of two phenomena: the precipitation of Ni during the crystallization processes and the change of the Ni saturation magnetization with temperature. Fig. 2 shows the magnetization of four Ni–P samples as a function of temperature. The corresponding DSC curves are also shown. The compositions of the four samples were chosen to be two hypoeutectic (16.9 and
362
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
Fig. 2. Magnetization and DSC curves of Ni – P samples taken at a heating rate of 3 K/min.
18.5 at.% P), one eutectic (19.2 at.% P) and one hypereutectic (19.9 at.% P). Generally, the change in magnetization occurs before any sign of heat flow. The
increase of P content results in a shift of the maxima of both types of curves, magnetothermal and DSC, to higher temperatures. The sample with 16.9 at.% P has markedly different magnetothermal curve, while its DSC curve is similar to that of the other two samples. The second step in the magnetothermal method is to calculate the ratio, marked as h(T), between magnetization during heating and cooling. If only FM particles are present in the sample then this ratio will determine the quantity of Ni precipitated at a given temperature as a fraction of the total quantity of Ni precipitated after complete crystallization [32]. The third step is to smooth and differentiate the h(T) ratio, in order to reveal the temperatures of maximum Ni precipitation rate. The differential curve dh(T)/dT is comparable with the DSC curve. The couples of h(T), dh(T)/dT and DSC curves for the hypoeutectic and eutectic alloys are presented in Fig. 3. Obviously, the major part of Ni
Fig. 3. The ratio h(T), its first derivative dh(T)/dT, microhardness and the DSC curve for (a) Ni83.1P16.9 and (b) Ni80.8P19.2 alloys.
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
363
3.3. Magnetothermal measurements of Ni– Me –P alloys
Fig. 4. Magnetization as a function of temperature for amorphous Ni–P, Ni– Sn– P and Ni– Cu–P alloys at a heating rate of 3 K/min.
precipitates at a temperature higher than 570 K since there is a major increase of h(T). Four peaks at 488, 543, 608 and 667 K are observed in the differentiated h(T) curve of the hypoeutectic sample with 16.9 at.% P (Fig. 3). The differentiated h(T) curve of the eutectic sample reveals three partially overlapping peaks at 558, 608 and 658 K. These features cannot be guessed from the magnetization curve. The microhardness of one hypoeutectic (16.4 at.% P) sample and one eutectic sample as a function of the end temperature is presented also in Fig. 3. Both samples show increase of the microhardness. Two areas may be noticed in the curves: one of slow growth and one of sharp increase. The slow growth area is from 450 to about 550 K for the hypoeutectic sample and from 490 to 570 K for the eutectic sample. The sharp increase area follows for both samples.
3.3.1. Ni –Sn –P Samples with a third element have lower magnetization in the whole temperature range, as seen in Fig. 4. The increase in Sn content also leads to a shift of the magnetization peak to higher temperature, especially for the MQ samples. The Curie temperature, TC, was determined by the linear fit of the reversed susceptibility of each alloy after complete crystallization. The values of Curie temperature are presented in Fig. 5 and Table 1 as a function of the Sn content. It may be seen that despite the rather inaccurate method of determination of the Curie temperature the linear fit is pretty good. 3.3.2. Ni –Cu –P As a rule, the addition of Cu reduces the magnetization more strongly than the addition of Sn. The Ni– Cu – P alloys show only a minute change in magnetization during heating (Fig. 4). A small increase is present at 570–580 K. The curves of the samples with different Cu content almost coincide. The Curie temperatures of the crystallized samples are also included in Table 1.
4. Discussion
4.1. Ni –P alloys The part of the Ni –P equilibrium phase diagram covering the range 16–20 at.% P is of eutectic type with terminal phases fcc Ni and bct Ni3P compound. The amorphous Ni –P alloy is thermodynamically unstable and upon heating it will tend to transform to a mixture
Fig. 5. The Curie temperature of the Ni –Sn–P alloys after complete crystallization as a function of the Sn content.
364
Sample
P (at.%)
Me (at.%)
Maximal possible quantity of Me in the Ni particles (at.%)
TC (K)
Quantity of Me in the particles according to TC (at.%)
Fraction of the Me in the particles (%)
Ni–Cu–P
9-11 9-12
19.8 9 0.3 20.3 9 0.3
2.2 90.2 2.8 90.2
10.6 15
525 9 20 521 917
10.9 11.3
100 75
Ni–Sn–P
ED-3 MQ-1 ED-4 MQ-4
199 1 18.0 9 0.4 209 1 17.9 9 0.6
0.29 0.1 0.40 9 0.04 0.79 0.2 0.90 9 0.04
0.8 1.4 3.5 3.1
605 9 23 588 9 22 568 9 36 553 929
0.8 1.2 2.0 2.3
100 86 57 74
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
Table 1 Content of the third element (Cu or Sn) in the sample and in the precipitated Ni particles
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
of crystalline Ni and Ni3P. The amorphous Ni – P alloy is paramagnetic. The Ni3P and all other Ni– P compounds are paramagnetic too. Thus, the FM properties of the crystalline Ni– P alloy are due only to the presence of fcc Ni particles. The maximum volume fraction of this phase after full crystallization is between 20 and 32% for the P content between 16 and 20 at.% and one should expect randomly distributed Ni particles in a paramagnetic matrix, either amorphous or crystalline Ni3P. The type of magnetic behavior of the Ni particles (as shown in Fig. 1 for room temperature) determines the magnetic properties of completely or partially crystalline Ni– P alloys. The magnetization versus temperature curve of our samples fits well to the crystallization behavior of the Ni –P alloy [6,9–11,18 – 20]. As seen in Fig. 2, the precipitation of the magnetic phase moves to higher temperatures when the P content is increased. Simultaneously, the quantity of the precipitated Ni diminishes. This is seen for samples with 16.9, 18.5 and 19.2 at.% P. With further increase of P content beyond the eutectic composition (19 at.%) a reaction of precipitation of Ni3P, Ni5P2 and Ni occurs at temperatures lower than that of the eutectic transformation in the eutectic alloy [9 –11]. That is the case with our samples too (Fig. 2). The steep increase of magnetization of the sample with 19.9 at.% P is at 520 K, while for the sample with 19.2 at.% P it is at 553 K. Note that the DSC peak of 19.9 at.% P sample is, however, at temperature higher than that of 19.2 at.% P sample. This means that for a sample with 19.9 at.% P the precipitation of the magnetic phase starts even before bulk crystallization occurs. Generally, it is clearly seen in Fig. 2 that the rise of magnetization for the eutectic and hypereutectic alloys occurs prior to the bulk crystallization (DSC peak). There is a difference of about 40– 70 K in the peak positions in magnetization versus temperature curves and the corresponding DSC peaks presented in Fig. 2. It is observed for both eutectic and hypereutectic samples and was also noted in our previous work [32] including two hypoeutectic MQ samples. The variation in the heating rate does not seem to affect this peak position difference [32]. A similar event may be found in Ref. [15], where the second peak in the magnetization curve of a hypoeutectic alloy and the first one of a eutectic alloy appear 80 K before the major DSC peaks. Lu et al. [26] investigated the crystallization of the MQ Ni80P20 alloy and observed only about 5 K difference between the peak in the susceptibility versus temperature curve and DSC curve taken at the same constant heating rate. In similar experiments of Wachtel et al. [11] a difference of 15 K was determined in peak position of susceptibility versus T and DSC curves, but the heating rates were 5 and 20 K/min, respectively.
365
The peak at 488 K in the dh(T)/dT curve of the hypoeutectic alloy (in Fig. 3) results from primary Ni crystallization reaction. Probably because of low scanning rate DSC is not discerning this reaction, whereas the magnetothermal method detects it very well. The next new moment is that the DSC peak denoting the bulk crystallization is very narrow, whereas the dh(T)/dT curve around it is very wide. In this wide temperature interval there is predominantly a change of the amount of the precipitated strong magnetic phase. The peaks at 543 K in the dh(T)/dT curve of the hypoeutectic sample and at 558 K in the curve of the eutectic sample correspond to an additional precipitation of Ni. It is more pronounced in the hypoeutectic alloy than in the eutectic one. The additional precipitation may be considered as a precursor of the eutectic reaction. It was stated in Ref. [1]: ‘‘the hypoeutectoid amorphous alloys does not have a driving force for Ni3P formation until the P concentration has increased beyond the amorph-Ni3P metastable equilibrium concentration by precipitating crystalline Ni.’’ In addition, it was established that the crystallization of the amorphous Ni–P is not the eutectic one and ‘‘the most stable amorphous phase against any changes in structure is in the range of 20–22 at.% P’’ [9]. If this composition will be considered as a quasieutectic one, then the alloy from Fig. 3b is quasihypoeutectic and we could apply the abovementioned explanations for the appearance of the hint of a small peak at 558 K. The primary and secondary precipitation of Ni is proved in an independent way by microhardness measurements made after the same isochronal heating as DSC and magnetothermal measurements. The data are shown in Fig. 3. Both samples have a temperature range in which the microhardness increases slowly. For the hypoeutectic sample this range is starting and finishing at lower temperatures (450–550 K), than for the eutectic alloy (490–570 K). In this range, the rise of the microhardness is due mainly to the precipitation of crystalline Ni particles in the amorphous matrix. The volume fraction of these particles is about 10% at 550 K for the hypoeutectic sample and about 2% at 570 K for the eutectic alloy as calculated from the h(T) curves. One could expect they are nanocrystalline. Besides Ni particles before the bulk crystallization one should expect a precipitation of Ni3(P,Ni). Limited formation of Ni3P was observed at temperatures below the DSC peak in ED Ni–3.6 at.% P [1]. The role of precursory Ni3P clusters for the crystallization micromechanism in the MQ Ni–P amorphous alloys has been discussed [39]. The steep increase in the microhardness curve could be connected with the formation of Ni3P particles in addition to Ni particles. The wellinvestigated phenomenon of precipitation hardening of monocrystalline materials is characterized with linear dependence on the reciprocal radius of precipitates and
366
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
with much weaker dependence on their volume fraction. (Relatively uniform amorphous matrix could be considered as closer to monocrystalline structure, than to the polycrystalline one with a strong influence of the grain boundaries.) We did not obtain information about the mechanism of the precipitation hardening in amorphous matrix, nevertheless a similar effect was observed in amorphous alloys [40]. When the bulk crystallization takes place, the volume fraction of the crystalline Ni3(P,Ni) phase increases more than three times faster than the volume fraction of Ni. The dependence of the microhardness of the Ni –P alloys on the average particle size has a maximum [41]. The hardness of nanocrystalline Ni80P20 alloy increases from 10 to 12 GPa when the average grain size grows from 10 to 100 nm [2]. Above 70– 100 nm the microhardness drops according to the Hall–Petch equation [42]. Therefore the highest microhardness is achieved in a completely crystalline sample with average grain size about 100 nm. The peak value of 12 GPa is obtained for our hypoeutectic sample after heating to 590 K. Coming back to the magnetothermal curve it looks like a good correlation exists between DSC peaks and magnetothermal curves maxima. As it is well known, the DSC peak reflects the bulk crystallization. The problem is that the DSC peaks are shifting to higher values, when the rate of isochronal heating is increased, but magnetothermal maxima around 600 K remain at the same place [32]. One should remember, that Ni Curie temperature (627 K) is crossed while measuring both the magnetization curves: during heating and during cooling. The strong para-process in the vicinity of the Curie temperature, as well as the abrupt change of the magnetization may considerably influence the h(T) and dh(T)/dT curves. Thus, both the height and position of the most prominent peak (at 608 K in Fig. 3) in dh(T)/dT curve could be affected. We should also note that beyond the Curie temperature the estimation of the Ni precipitation and dh(T)/dT is correct, only if the paramagnetic susceptibility of Ni is several times greater than that of Ni3P. There are indications that this is true [32]. The last peak at 650– 670 K is in the paramagnetic range. However, it is not present in the MQ samples [32]. Therefore, we may consider it as a consequence of some inhomogeneity or peculiarity of the ED samples. Thus the utilization of the magnetothermal method in combination with DSC and microhardness measurements reveals new details of the magnetic phase precipitation.
4.2. The addition of a third element The addition of a third element can affect the magnetization versus temperature curve in two ways. The first
is by delaying the precipitation process. This will result in a shift of the peak to higher temperatures, when compared with the peak of an alloy with the same content of P, but without a third element. The second way is by affecting the phase to be precipitated and particularly its magnetic properties, rather than the kinetics of the process. This case for Ni–Me –P with given contents will result in changing the peak height in the magnetization versus temperature curve. Both effects are present for our samples. The third element can either be dissolved in the fcc Ni phase yielding Ni(Me) solid solution, or in the Ni3P phase or may form compounds with P. To the XRD sensitivity Cu compounds with P (such as Cu3P) were not found in Ni–Cu –P system with up to 50% Cu [21]. No data are available for the appearance of Sn–P or Ni –Sn compounds in the Ni–Sn –P system at low content of Sn. None of these compounds which might be formed during sample crystallization is FM. Thus, the only phase with strong magnetic properties that may be present in the sample, is the Ni or Ni(Me,P) solid solution. The two metals, Cu and Sn, are related differently to Ni. Copper is unlimitedly soluble in Ni, while Sn has only very limited solubility in Ni [43]. However, tin has four times stronger effect on the magnetic moment of Ni, than Cu has.
4.2.1. Addition of Sn As a whole, the magnetization versus temperature curve of the Ni–Sn–P alloys is similar to that of the Ni –P alloys (Fig. 4). If one accounts for the position and height of the peaks, then both effects of the third element discussed above are seen: the retardation of the transformation and the reduction of the magnetization due to the alloying of the precipitated Ni particles with Sn. Let us consider the MQ Ni–Sn –P alloys in Fig. 4. The precipitation process in the 0.40 at.% Sn sample begins at almost the same temperature like in the sample with 19 at.% P, not containing Sn. However, the speed of magnetization increase is not as high as for the sample without Sn and the maximum magnetization achieved is lower. The MQ sample with 0.90 at.% Sn has the highest temperature of steepest magnetization increase and further diminution in the height of the magnetization maximum. Keeping in mind that the sample without Sn contains 19 at.% P, one should suppose that the sample with 18 at.% would have even lower temperature of steepest magnetization increase. Thus, in the MQ samples, the addition of Sn retards the Ni precipitation process and therefore the crystallization of the sample. The last statements are obviously true also for the ED samples. However, the increase of both Sn content from 0.2 to 0.7 at.% and P content from 19 to 20 at.%, does not lead to an increase of the
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
367
Ni-P Fig. 6. Relative magnetization h(T) =| Ni-Me-P heating / | cooling curves of amorphous alloys heated at a rate of 3 K/min.
temperature of steepest magnetization increase. This may be due to the opposite effects of the increase of both phosphorus (see Figs. 2 and 4) and tin. The retardation of the crystallization processes due to the introduction of Sn in Ni– P observed in this work using magnetic measurements is in correlation with the earlier DSC studies [18]. A shift of the DSC peak of the ED Ni –Sn –P alloys to the higher temperature was found in comparison with its position in the case of the ED Ni – P alloy. For the ternary Ni– Me – P alloys the precipitating FM phase is not pure Ni. Therefore, the saturation magnetization varies not only with the temperature and phase quantity, but also with the content of the third element in the precipitating phase. This is true for both heating and cooling the sample. Samples with different content of third element have differing cooling curves. This renders their h(T) ratios incomparable. Thus, it is more beneficial to divide by magnetization during heat-
ing to the temperature dependence of the magnetization of pure Ni, rather than by the magnetization of the sample during cooling. Before the division, the Ni magnetization versus temperature curve should be rescaled to the quantity of pure Ni that would precipitate after complete crystallization as if the alloy was Ni–P. Then h(T) will be the ratio between the magnetic moment of the precipitating FM phase and the magnetic moment the sample would have after complete crystallization if pure Ni from Ni–P were precipitating. As can be seen from Fig. 6, the h(T) curve of the Ni –Sn–P alloys is similar to those of the Ni–P alloys, but does not tend to 1 at higher temperatures. This is due to the presence of Sn in the precipitating Ni phase. Moreover, there is a decrease above 670 K for samples with higher Sn content. This manifests in the negative values of dh(T)/dT curves in Fig. 7 and indicates that Sn incorporates in the particle to a larger extent at higher temperatures.
Ni-P Fig. 7. Differentiated h(T):dh(T) / dT= d(| Ni-Me-P heating / | cooling) / dT curves of amorphous alloys heated at a rate of 3 K/min.
368
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
As a whole, the dh(T)/dT curves of the Ni–Sn –P alloys are also similar to that of the Ni– P alloy. They appear at the same position and have similar shape, but the distinct peaks in the Ni– P alloy seem to be ‘smoothened’ in the Ni–Sn – P alloys. Thus, the addition of Sn retards the precipitation of magnetic phase, which is now Ni(Sn) instead of Ni. The lower maxima are evidence for lower rate of magnetization increase. It is also clear from Fig. 7 that the precipitating FM phase contains Sn in a quantity proportional to the Sn content in the sample. This phase is most probably a supersaturated solution of Sn in Ni, and the content of Sn in this phase varies in the course of the phase transformation.
4.2.2. The addition of Cu The addition of Cu has an even more dramatic effect on the magnetization versus temperature curve, i.e. the magnetization slightly changes from the level of the paramagnetic amorphous phase. However, the small increase in the range 570– 580 K indicates a common behavior with a high P/Sn samples, i.e. Cu retards the crystallization no more than Sn does, as was already shown by DSC measurements [18]. Copper incorporates into the precipitating Ni particles to a quantity that is at least four times larger. Table 1 shows that this is true for the fully crystallized samples. The reason for this is the higher solubility of Cu in Ni. Since there is very small change in the magnetization of the Ni –Cu – P alloys during heating the h(T) and dh(T)/dT curves were not calculated. The Ni– Cu –P alloys retained lowest magnetization also after complete crystallization and cooling to room temperature, so that the copper is a better additive if strong FM properties have to be reduced. 4.3. The magnetization during cooling From the Curie temperatures estimated through the magnetization during cooling after the annealing step at 800 K the content of Sn and Cu in the precipitated Ni particles was evaluated and is presented in Table 1. The solubility of Sn in Ni at the annealing temperature, 800 K, is about 1 at.%. For three of four samples the content of Sn in the Ni particles exceeds the equilibrium solubility limit. Thus, half an hour at 800 K was not enough for the alloys to reach the equilibrium state. Therefore the estimated Curie temperatures and the magnetization during cooling can be used to draw some conclusion about the crystallization process. The incorporation of Sn in the precipitating Ni particles seems to be proportional to the Sn content in the entire sample. Also it does not depend on the type of sample preparation: MQ or ED. The same cannot be said for the shape of the magnetization versus temperature curve during cooling. The shape of the magnetiza-
tion curve of the MQ samples is very close to that of the electroless Ni– P samples and pure Ni while the shape of the electroless Ni–Sn– P samples is slightly different. This makes us think that the addition of Sn in the ED Ni –P samples affects the kinetics of Ni particles precipitation by a way different from that of the MQ samples, which results in different final structure.
5. Conclusion The crystallization of amorphous Ni–P, Ni –Sn–P and Ni– Cu –P alloys is investigated by a modification of a magnetic phase analysis method called magnetothermal method. It consists of three steps. The first step is the measurement of magnetization during heating. Then annealing to complete crystallization of the sample is conducted and the magnetization during cooling is measured. Then the ratio between two magnetization values (during the heating and cooling) is constructed as dependent on temperature, h(T), which is the second step. The third step is the calculation of the derivative: dh(T)/dT. The last operation is making the magnetothermal curve directly comparable to DSC curves, as far as peak positions are concerned. The magnetothermal method is much more sensitive to the precipitation of Ni or Ni-based phases during isochronal heating than DSC and this is demonstrated at a heating rate of 3 K/min. In this way, the details of the precipitation processes could be investigated. Particularly, it was shown that in the hypoeutectic alloys besides the expected primary Ni precipitation, there is an additional secondary Ni precipitation. This secondary precipitation could be noticed to a smaller extent in eutectic and even in the hypereutectic alloys with phosphorus content close to the eutectic composition. The microhardness measurement during isochronal heating also reveals the Ni precipitation in amorphous Ni –P matrix. The effect of this precipitation on microhardness, as well as estimation based on magnetic measurements suggests that the precipitated particles are of nanometer size. The addition of a third element to the Ni–P amorphous alloy can affect its thermomagnetic behavior in two ways. First, by delaying the strong magnetic phase precipitation process. Second, by affecting the magnetic properties of the precipitating phase due to alloying. The addition of Sn or Cu to the Ni–P alloy leads to the incorporation of the third element into the precipitating Ni particles. Since both Sn and Cu decrease the spontaneous magnetization of Ni, they diminish the magnetization of the Ni–Me –P alloy. Simultaneously they retard the precipitation of the Ni(Me) phase. Though Sn could have a stronger effect on the Ni spontaneous magnetization, it decreases the magnetization of the
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
Ni –P alloy to a lesser extent than Cu does, due to the lesser solubility of Sn in Ni. However, the addition of Sn shifts the crystallization process to higher temperatures. Thus, the addition of Cu or Sn reduces the appearance of phase with strong magnetic properties. Obviously, the proposed magnetothermal method in combination with DSC and microhardness measurements could give valuable information for the kinetics of crystallization of multicomponent amorphous alloys containing a strong magnetic phase. In this way, the use of magnetic measurements becomes much more helpful in investigating the formation of magnetic nanoparticles in an amorphous para- or diamagnetic matrix, even though the interpretation of magnetization data is not always easy and straightforward.
[12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24]
Acknowledgements
[25]
The authors are indebted to Professor Dr Sc. I. Avramov from the Institute of Physical Chemistry of Bulgarian Academy of Sciences for the DSC measurements, the Copernicus program for the DSC equipment and to Mr Ralev from DZU Plc. in Stara Zagora for access to the VSM.
[26] [27] [28] [29] [30] [31] [32] [33] [34] [35]
References [1] Ts. Hentschel, D. Isheim, R. Kirchheim, F. Mueller, H. Kreye, Acta Mater. 48 (2000) 933. [2] K. Lu, Mater. Sci. Eng. R16 (1996) 161. [3] R.N. Duncan, Paper presented at COATINGS’99, Conference Proceedings, Dallas, 21 –23 September, 1999, p. 265. [4] J. Hajdu, Trans. Inst. Met. Finishing 75 (1) (1997) B7. [5] D.R. Frear, J. Met. 51 (3) (1999) 22. [6] U. Pittermann, S. Ripper, Phys. Stat. Sol. (a) 93 (1986) 131. [7] U. Pittermann, S. Ripper, in: S. Steeb, H. Warlimont (Eds.), Rapidly Quenched Metals, Elsevier, Amsterdam, 1985, p. 375. [8] U. Pittermann, S. Ripper, Z. Metall. 74 (1983) 783. [9] K.-H. Hur, J.-H. Jeong, D.N. Lee, J. Mater. Sci. 25 (1990) 2573. [10] K.-H. Hur, J.-H. Jeong, D.N. Lee, in: L. Huang (Ed.), Thin Films and Beam-Solid Inter, Elsevier, Amsterdam, 1991, p. 247. [11] E. Wachtel, I. Bakonyi, J. Bahle, N. Willmann, A. Burgstaller, W. Socher, J. Voitla¨ nder, A. Lovas, H.H. Libermann, Mater. Sci. Eng. A 133 (1991) 196.
.
[36] [37] [38] [39] [40] [41]
[42]
[43]
369
K. Lu, J.T. Wang, J. Non-Cryst. Solids 117/118 (1990) 716. K. Lu, J.T. Wang, Mater. Sci. Eng. A 133 (1991) 500. K. Lu, J.T. Wang, J. Cryst. Growth 94 (1989) 448. K. Masui, M. Tachihara, T. Yamada, T. Tsujimoto, J. Jpn. Inst. Met. 44 (1980) 124. K. Masui, S. Maruno, T. Yamada, J. Jpn. Inst. Met. 41 (1977) 1130. K. Masui, T. Yamada, Y. Hisamatsu, J. Met. Fin. Soc. Jpn. 31 (1980) 667. N. Krasteva, S. Armyanov, J. Georgieva, N. Avramova, V. Fotty, J. Electron. Mater. 24 (1995) 941. N. Krasteva, V. Fotty, S. Armyanov, J. Electochem. Soc. 141 (1994) 2864. S. Armyanov, N. Krasteva, V. Fotty, Proc. Int. Symp. GALVANO’91, Varna, 30 September – 2 October 1991, p. 240. K.-H. Hur, J.-H. Jeong, D.N. Lee, J. Mater. Sci. 26 (1991) 2037. X. Haowen, Zh. Bangwei, Y. Qiaoqin, Trans. IMF 77 (1999) 99. B.A. Apaev, Fazoviy Magnitniy Analiz Splavov, Metallurgia, Moskva, 1976 (in Russian). A. Berkowitz, E. Kneller (Eds.), Magnetism and Metallurgy, vols. 1 – 2, Academic Press, New York, 1969. E. Wachtel, N. Willmann, B. Predel, J. Magn. Magn. Mater. 59 (1986) 115. K. Lu, R. Lu¨ ck, B. Predel, J. Alloys Comp. 201 (1993) 229. R. Lu¨ ck, J. Mater. Sci. Technol. 13 (1997) 260. R. Lu¨ ck, K. Lu, W. Frantz, Scr. Metall. Mater. 28 (1993) 1071. A. Aronin, G. Abrosimova, I. Zver’kova, D. Lang, R. Lu¨ ck, J. Non-Cryst. Solids 208 (1996) 139. H.G. Zolla, F. Spaepen, Mater. Sci. Eng. A 204 (1995) 71. H.G. Zolla, F. Spaepen, IEEE Trans. Magn. 31 (1995) 3814. D. Tachev, D. Iorgov, S. Armyanov, J. Non-Cryst. Solids 270 (2000) 66. A. Amamou, J. Durand, Commun. Phys. 1 (1976) 191. J.S. Kouvel, J.B. Comly, Phys. Rev. Lett. 24 (1970) 598. I. Baconyi, A. Burgstaller, W. Socher, J. Voitla¨ nder, E. TothKadar, A. Lovas, H. Ebert, E. Wachtel, N. Willmann, H.H. Liebermann, Phys. Rev. B 47 (1993) 14961. A. Aharoni, IEEE Trans. Magn. 29 (1993) 2596. W.F. Brown, Ann. N. Y. Acad. Sci. 147 (1969) 463. E.C. Stoner, E.P. Wolhfarth, Phil. Trans. R. Soc. A 240 (1948) 599. K. Lu, J.T. Wang, Mater. Sci. Eng. A 133 (1991) 504. H. Kimura, M. Masumoto, in: F.E. Luborsky (Ed.), Amorphous Metallic Alloys, Butterworths, London, 1983 (chap. 11). R. Rofagha, U. Erb, in: H.D. Merchant (Ed.), Defect Structure, Morphology and Properties of Deposits, The Minerals, Metals and Materials Society, 1995, p. 245. Y. Tobiyama, B. Derby, J.B. Pethica, in: H.D. Merchant (Ed.), Defect Structure, Morphology and Properties of Deposits, The Minerals, Metals and Materials Society, 1995, p. 335. T.B. Massalski (Editor-in-Chief), Binary Alloy Phase Diagrams, vols. 1 – 2, ASM, Ohio, 1986.
Magnetothermal study of nanocrystalline particle formation in amorphous electroless Ni–P and Ni–Me–P alloys D. Tachev, J. Georgieva, S. Armyanov * Institute of Physical Chemistry, Bulgarian Academy of Sciences, Sofia 1113, Bulgaria Received 1 October 2000; received in revised form 13 March 2001
Abstract The microstructure of amorphous Ni–P and Ni–Me– P materials and especially its change during the heat treatment is of great importance for their magnetic, mechanical and corrosion behavior. A new magnetic phase analysis method (magnetothermal) is presented that reveals the precipitation of nanoparticles with strong magnetic properties during phase transformation upon heat treatment. It is applied to electroless Ni–P, Ni–Cu–P and Ni– Sn – P amorphous alloys. The results acquired by this method are compared with data obtained by differential scanning calorimetry, as well as by microhardness measurements using identical heat treatment in all three cases. Due to the high sensitivity of the magnetothermal method a more detailed picture of the precipitation processes in Ni–P alloys is obtained and the new information is discussed. Magnetothermal measurements reveal several stages of precipitation of a phase with strong magnetic properties. This phase is Ni in the Ni– P alloy, and Ni(Me) solid solution in the Ni–Me–P alloys. Though Sn has a stronger effect on the Ni magnetization, Cu is more effective in preventing the appearance of high magnetization in a thermally treated Ni–Cu–P alloy. This is due to Cu incorporation in Ni particles in a quantity above four times larger than Sn. © 2001 Elsevier Science Ltd. All rights reserved. Keywords: Ni– P, Ni– Cu–P and Ni–Sn–P alloys; Crystallization; Magnetic properties; Microhardness; Differential scanning calorimetry
1. Introduction Electrolessly deposited (ED) Ni – P alloys are unique engineering materials. With the now growing interest in nanomaterials the ED Ni – P alloys remain in focus. Low phosphorus alloys are nanocrystalline as plated [1], whereas high phosphorus alloys are brought to nanocrystalline state by crystallizing an amorphous alloy [2]. Ni –P coatings have a bright and shiny appearance, outstanding mechanical properties and they adhere strongly to the substrate. They possess excellent wear and corrosion resistance. The ED Ni – P alloys with different composition find applications in many branches of industry from food and textile to automotive and aircraft [3]. For example, ED Ni – P utilization in thin film disks is well known [4] and its possible use in flip-chip packaging is currently under investigation [5]. * Corresponding author. E-mail address: [email protected] (S. Armyanov).
The crystallization of Ni –P amorphous alloys has long been studied by electron diffraction, X-ray diffraction (XRD), transmission electron microscopy and differential scanning calorimetry (DSC) combined with other methods including microhardness and magnetic measurements [6–17]. The overall crystallization process can be summarized as follows [9–11,18 –20]: for hypoeutectic alloys (in the range 15 –19 at.% P) a reaction of the primary crystallization of Ni is observed at low temperature and a eutectic transformation, with Ni and Ni3P as final products, at high temperature. Only a eutectic reaction occurs in alloys with eutectic composition (19 –20 at.% P). In hypereutectic alloys, Ni, Ni3P and Ni5P2 crystalline phases form at low temperature. At higher temperature the metastable Ni5P2 phase transforms into Ni3P. All these reactions are present as distinct peaks in the DSC thermograms. The peak position dependence on phosphorus content is summarized in Refs. [9–11]. Ternary alloys, Ni –Me –P, based on the Ni –P binary system were introduced because of the necessity of increasing the thermal stability of the paramagnetic
0013-4686/01/$ - see front matter © 2001 Elsevier Science Ltd. All rights reserved. PII: S 0 0 1 3 - 4 6 8 6 ( 0 1 ) 0 0 5 8 7 - 4
360
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
properties of amorphous Ni– P alloys used as an underlayer in thin film memory disks [19,21]. Nontransition elements have been added to the Ni– P alloys as a third component namely Cu [10,18– 21], Sn [18,22] and Sb [18]. It was shown that Ni– Cu alloy’s particles form in the Ni – Cu –P alloys and this decreases their magnetization after heat treatment [18,19,21]. Both Cu and Sn increase the crystallization temperature and the crystallization activation energy and diminish the alloy magnetization after heat treatment [18]. The highest thermal stability of the amorphous state was achieved by the addition of Sb, but the magnetization of the annealed samples was also high [18]. Magnetic measurements are generally used when a certain magnetic property is the target or for investigating the origin of certain magnetic behavior. Despite the fact that magnetic phase analysis has been known for a long time [23,24] magnetic measurements are rarely used for the characterization of sample structure and phase transitions. This is due to the difficulty in obtaining quantitative information. Magnetic phase analysis is based on the fact that the magnetic moment of a sample consisting of several phases is a sum of their moments and each phase magnetization has characteristic temperature dependence. The aim is to find the mass or volume fraction of each phase. The method is applicable to phases with strong magnetic properties. Masui et al. [15] performed magnetization measurements during the heating of amorphous Ni– P samples. Since the sample is initially paramagnetic and Ni precipitates during crystallization, the stages of crystallization were traced. Similar experiments were made by measuring magnetic susceptibility instead of magnetization [25–29]. Lu¨ ck et al. [27,28] applied the Johnson– Mehl –Kolmogorov – Avrami analysis to the susceptibility curves. However, their assumption [28] that the susceptibility is proportional to the volume of the crystalline phases is rather presumptuous, since the magnetic response of nanometer size Ni solid solution particles is essentially nonlinear. A better approach to the investigation of small particles of strong magnetic phases in a para- or diamagnetic matrix is superparamagnetic (SPM) granulometry [30,31]. It is capable of finding magnetic particles’ size distribution, as well as their volume fraction. The basic property behind magnetization and susceptibility is the magnetic moment. It is a function of both the magnetic phase volume and composition. This is a significant difficulty especially for studies of phase transitions involving simultaneous change of the volume and composition of a phase with strong magnetic properties. Determination of the chemical composition by an independent method is required for phase identification. Magnetic measurements, however, have a number of advantages. First, we should mention the sensitivity.
Magnetization measurement with a standard vibrating sample magnetometer (VSM) could be up to hundred times more sensitive to the volume change of a Ni phase in the Ni–P sample than DSC. Thus the magnetization measurements are a better choice for early detection of a transformation involving phases with strong magnetic properties [28]. Second, the magnetic response is selective. It is very large for phases with strong magnetic properties compared to phases with weak magnetic properties. Thus, only the formation and disappearance of phases with strong magnetic properties will be discerned contrary to DSC, that will detect any transformation connected with substantial heat release. Third, the magnetic response is practically immediate. This is particularly important for the investigation of kinetics. In order to receive more comprehensive information for the crystallization of amorphous Ni–P besides the abovementioned combination of several techniques, the necessity of introducing new methods is obvious. Recently, the results of the use of atom-probe field ion microscopy in combination with DSC and XRD have been summarized [1]. In our previous work [32], we described a new magnetothermal method appropriate for approximate determination of the total volume of the SPM and ferromagnetic (FM) particles precipitating in amorphous Ni–P alloys. In this paper, we will use the magnetothermal method to reveal more details about the formation of nanocrystalline particles in the course of the crystallization process of the Ni–P, Ni – Sn–P and Ni –Cu –P alloys. The results will be compared with DSC and microhardness measurements applying identical heat treatment.
2. Experimental The samples used in this investigation were amorphous ED foils of the Ni1 − x Px alloy (after dissolution of the substrate), as well as melt-quenched (MQ) ribbons and ED foils of the Ni1 − x − y Mey Px alloy, with 16B xB 20 and 0.2ByB 3.0. Electroless (autocatalytical) deposition was performed in Ni sulfate solution, using hypophosphite as the reducing agent. The chemical composition of the plated alloys was determined using energy dispersive spectrometer on a scanning electron microscope JEOL 733. The changes of magnetization at field strengths of 8– 12 kOe were measured while heating the sample from 420 to 720 K at a constant rate of 3 K/min (isochronal heating). Samples were subsequently heated faster to 800 K and held there for about half an hour to completely crystallize them. The magnetization was also recorded during subsequent cooling at 12 K/min. A second cycle of heating and cooling was performed to assure that no further crystallization and accompanying
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
change in magnetization were present. This confirmed that the samples crystallized completely for half an hour at 800 K. All measurements were performed on a VSM, Riken Denshi BHV-55, under vacuum better than 10 − 4 Torr. The DSC measurements were performed in a Perkin–Elmer DSC-7 using N2 as a purge gas. The mass of each sample was 3– 12 mg. The scan rates were identical to those of the magnetothermal measurements. Samples for microhardness measurement were heated with a constant heating rate of 3 K/min in a programmable oven. They were put in the oven at 390 K, pulled out at different temperatures and quenched in cold water. The microhardness was tested by the Vickers method. The load of 245.2 mN was applied for 20 s. To ensure that the thickness of our samples was enough for the correct hardness measurement with this load the softest samples were measured also at a load of 147.1 mN. The same microhardness value was obtained.
3. Results
3.1. Magnetic beha6ior of Ni particles depending on their size The Ni – P alloys represent most accessible case for magnetic phase analysis of precipitation of particles of strong magnetic phase in a matrix with weak magnetic phase properties. The magnetic behavior of pure Ni particles strongly depends on their size and systematization of this behavior based on the data from the literature is presented here. Four groups could be considered. Starting from the smallest size they are giant moment paramagnetic (GMP) clusters, SPM, single domain (SD) and FM particles (Fig. 1). Coupling between atomic magnetic moments, that is characteristic for strong magnetic phase, exists in all four types of particles. The GMP clusters consist of a small number of atoms. The magnetic moment of the cluster is a sum of the moments of the atoms. According to quantum mechanics the magnitude and the projection of this moment can take only discrete values. Ni particles with linear dimensions less than 4 A, behave like GMP
Fig. 1. Size dependence of the magnetic behavior of isolated (magnetically non-interacting) Ni particles at 300 K.
361
clusters [33,34], while particles of 7–8 A, are described as SPM [35]. Due to exchange interaction the SPM particles have significant magnetic moment, but their size is still too small to keep it constant. Thus, because of the thermal agitation their magnetic moment changes its direction and magnitude, so that no remanent magnetization and coercive force can be observed. The magnetic moments of these particles behave like moments of atoms of a paramagnetic material. The upper limit for SPM behavior of Ni particles at 300 K is about 188 A, . With increasing particle size the thermal agitation can no longer disorder the atom moments in the particle, so a hysteresis loop is formed. However, the size of these particles is comparable to the magnetic domain wall thickness and domains cannot be formed. Thus, the particles are called SD particles. The SD particle is always magnetized to almost complete saturation [36]. Two modes of magnetization reversal in this type of particles are possible. A coherent rotation of the magnetization is observed, when all atomic moments in the particle rotate uniformly, and incoherent rotation occurs, when atomic moments do not remain parallel during magnetization reversal. With further increase of particle size its volume becomes large enough to contain more than one magnetic domain. The particle becomes multidomain like most bulk FM bodies. The rigorous mathematical consideration shows that a range of dimensions can mark the transition from SD to MD behavior, rather than a single value [37]. Despite this, the value of 382 A, is given for Ni particles. From equations presented in Ref. [38] the range of 312– 1260 A, can be estimated for Ni at 300 K. It must be noted that all boundary dimensions are temperature dependent. Fig. 1 gives the size dependence of the magnetic behavior of Ni particles at 300 K. Mostly SPM particles are encountered when investigating the crystallization of the Ni– P amorphous alloys.
3.2. Magnetothermal, DSC and microhardness measurements of Ni–P alloy with different compositions The first step in the magnetothermal method is to measure the magnetization at constant magnetizing field strength during isochronal heating. Then the change of magnetization is a result of the combined action of two phenomena: the precipitation of Ni during the crystallization processes and the change of the Ni saturation magnetization with temperature. Fig. 2 shows the magnetization of four Ni–P samples as a function of temperature. The corresponding DSC curves are also shown. The compositions of the four samples were chosen to be two hypoeutectic (16.9 and
362
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
Fig. 2. Magnetization and DSC curves of Ni – P samples taken at a heating rate of 3 K/min.
18.5 at.% P), one eutectic (19.2 at.% P) and one hypereutectic (19.9 at.% P). Generally, the change in magnetization occurs before any sign of heat flow. The
increase of P content results in a shift of the maxima of both types of curves, magnetothermal and DSC, to higher temperatures. The sample with 16.9 at.% P has markedly different magnetothermal curve, while its DSC curve is similar to that of the other two samples. The second step in the magnetothermal method is to calculate the ratio, marked as h(T), between magnetization during heating and cooling. If only FM particles are present in the sample then this ratio will determine the quantity of Ni precipitated at a given temperature as a fraction of the total quantity of Ni precipitated after complete crystallization [32]. The third step is to smooth and differentiate the h(T) ratio, in order to reveal the temperatures of maximum Ni precipitation rate. The differential curve dh(T)/dT is comparable with the DSC curve. The couples of h(T), dh(T)/dT and DSC curves for the hypoeutectic and eutectic alloys are presented in Fig. 3. Obviously, the major part of Ni
Fig. 3. The ratio h(T), its first derivative dh(T)/dT, microhardness and the DSC curve for (a) Ni83.1P16.9 and (b) Ni80.8P19.2 alloys.
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
363
3.3. Magnetothermal measurements of Ni– Me –P alloys
Fig. 4. Magnetization as a function of temperature for amorphous Ni–P, Ni– Sn– P and Ni– Cu–P alloys at a heating rate of 3 K/min.
precipitates at a temperature higher than 570 K since there is a major increase of h(T). Four peaks at 488, 543, 608 and 667 K are observed in the differentiated h(T) curve of the hypoeutectic sample with 16.9 at.% P (Fig. 3). The differentiated h(T) curve of the eutectic sample reveals three partially overlapping peaks at 558, 608 and 658 K. These features cannot be guessed from the magnetization curve. The microhardness of one hypoeutectic (16.4 at.% P) sample and one eutectic sample as a function of the end temperature is presented also in Fig. 3. Both samples show increase of the microhardness. Two areas may be noticed in the curves: one of slow growth and one of sharp increase. The slow growth area is from 450 to about 550 K for the hypoeutectic sample and from 490 to 570 K for the eutectic sample. The sharp increase area follows for both samples.
3.3.1. Ni –Sn –P Samples with a third element have lower magnetization in the whole temperature range, as seen in Fig. 4. The increase in Sn content also leads to a shift of the magnetization peak to higher temperature, especially for the MQ samples. The Curie temperature, TC, was determined by the linear fit of the reversed susceptibility of each alloy after complete crystallization. The values of Curie temperature are presented in Fig. 5 and Table 1 as a function of the Sn content. It may be seen that despite the rather inaccurate method of determination of the Curie temperature the linear fit is pretty good. 3.3.2. Ni –Cu –P As a rule, the addition of Cu reduces the magnetization more strongly than the addition of Sn. The Ni– Cu – P alloys show only a minute change in magnetization during heating (Fig. 4). A small increase is present at 570–580 K. The curves of the samples with different Cu content almost coincide. The Curie temperatures of the crystallized samples are also included in Table 1.
4. Discussion
4.1. Ni –P alloys The part of the Ni –P equilibrium phase diagram covering the range 16–20 at.% P is of eutectic type with terminal phases fcc Ni and bct Ni3P compound. The amorphous Ni –P alloy is thermodynamically unstable and upon heating it will tend to transform to a mixture
Fig. 5. The Curie temperature of the Ni –Sn–P alloys after complete crystallization as a function of the Sn content.
364
Sample
P (at.%)
Me (at.%)
Maximal possible quantity of Me in the Ni particles (at.%)
TC (K)
Quantity of Me in the particles according to TC (at.%)
Fraction of the Me in the particles (%)
Ni–Cu–P
9-11 9-12
19.8 9 0.3 20.3 9 0.3
2.2 90.2 2.8 90.2
10.6 15
525 9 20 521 917
10.9 11.3
100 75
Ni–Sn–P
ED-3 MQ-1 ED-4 MQ-4
199 1 18.0 9 0.4 209 1 17.9 9 0.6
0.29 0.1 0.40 9 0.04 0.79 0.2 0.90 9 0.04
0.8 1.4 3.5 3.1
605 9 23 588 9 22 568 9 36 553 929
0.8 1.2 2.0 2.3
100 86 57 74
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
Table 1 Content of the third element (Cu or Sn) in the sample and in the precipitated Ni particles
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
of crystalline Ni and Ni3P. The amorphous Ni – P alloy is paramagnetic. The Ni3P and all other Ni– P compounds are paramagnetic too. Thus, the FM properties of the crystalline Ni– P alloy are due only to the presence of fcc Ni particles. The maximum volume fraction of this phase after full crystallization is between 20 and 32% for the P content between 16 and 20 at.% and one should expect randomly distributed Ni particles in a paramagnetic matrix, either amorphous or crystalline Ni3P. The type of magnetic behavior of the Ni particles (as shown in Fig. 1 for room temperature) determines the magnetic properties of completely or partially crystalline Ni– P alloys. The magnetization versus temperature curve of our samples fits well to the crystallization behavior of the Ni –P alloy [6,9–11,18 – 20]. As seen in Fig. 2, the precipitation of the magnetic phase moves to higher temperatures when the P content is increased. Simultaneously, the quantity of the precipitated Ni diminishes. This is seen for samples with 16.9, 18.5 and 19.2 at.% P. With further increase of P content beyond the eutectic composition (19 at.%) a reaction of precipitation of Ni3P, Ni5P2 and Ni occurs at temperatures lower than that of the eutectic transformation in the eutectic alloy [9 –11]. That is the case with our samples too (Fig. 2). The steep increase of magnetization of the sample with 19.9 at.% P is at 520 K, while for the sample with 19.2 at.% P it is at 553 K. Note that the DSC peak of 19.9 at.% P sample is, however, at temperature higher than that of 19.2 at.% P sample. This means that for a sample with 19.9 at.% P the precipitation of the magnetic phase starts even before bulk crystallization occurs. Generally, it is clearly seen in Fig. 2 that the rise of magnetization for the eutectic and hypereutectic alloys occurs prior to the bulk crystallization (DSC peak). There is a difference of about 40– 70 K in the peak positions in magnetization versus temperature curves and the corresponding DSC peaks presented in Fig. 2. It is observed for both eutectic and hypereutectic samples and was also noted in our previous work [32] including two hypoeutectic MQ samples. The variation in the heating rate does not seem to affect this peak position difference [32]. A similar event may be found in Ref. [15], where the second peak in the magnetization curve of a hypoeutectic alloy and the first one of a eutectic alloy appear 80 K before the major DSC peaks. Lu et al. [26] investigated the crystallization of the MQ Ni80P20 alloy and observed only about 5 K difference between the peak in the susceptibility versus temperature curve and DSC curve taken at the same constant heating rate. In similar experiments of Wachtel et al. [11] a difference of 15 K was determined in peak position of susceptibility versus T and DSC curves, but the heating rates were 5 and 20 K/min, respectively.
365
The peak at 488 K in the dh(T)/dT curve of the hypoeutectic alloy (in Fig. 3) results from primary Ni crystallization reaction. Probably because of low scanning rate DSC is not discerning this reaction, whereas the magnetothermal method detects it very well. The next new moment is that the DSC peak denoting the bulk crystallization is very narrow, whereas the dh(T)/dT curve around it is very wide. In this wide temperature interval there is predominantly a change of the amount of the precipitated strong magnetic phase. The peaks at 543 K in the dh(T)/dT curve of the hypoeutectic sample and at 558 K in the curve of the eutectic sample correspond to an additional precipitation of Ni. It is more pronounced in the hypoeutectic alloy than in the eutectic one. The additional precipitation may be considered as a precursor of the eutectic reaction. It was stated in Ref. [1]: ‘‘the hypoeutectoid amorphous alloys does not have a driving force for Ni3P formation until the P concentration has increased beyond the amorph-Ni3P metastable equilibrium concentration by precipitating crystalline Ni.’’ In addition, it was established that the crystallization of the amorphous Ni–P is not the eutectic one and ‘‘the most stable amorphous phase against any changes in structure is in the range of 20–22 at.% P’’ [9]. If this composition will be considered as a quasieutectic one, then the alloy from Fig. 3b is quasihypoeutectic and we could apply the abovementioned explanations for the appearance of the hint of a small peak at 558 K. The primary and secondary precipitation of Ni is proved in an independent way by microhardness measurements made after the same isochronal heating as DSC and magnetothermal measurements. The data are shown in Fig. 3. Both samples have a temperature range in which the microhardness increases slowly. For the hypoeutectic sample this range is starting and finishing at lower temperatures (450–550 K), than for the eutectic alloy (490–570 K). In this range, the rise of the microhardness is due mainly to the precipitation of crystalline Ni particles in the amorphous matrix. The volume fraction of these particles is about 10% at 550 K for the hypoeutectic sample and about 2% at 570 K for the eutectic alloy as calculated from the h(T) curves. One could expect they are nanocrystalline. Besides Ni particles before the bulk crystallization one should expect a precipitation of Ni3(P,Ni). Limited formation of Ni3P was observed at temperatures below the DSC peak in ED Ni–3.6 at.% P [1]. The role of precursory Ni3P clusters for the crystallization micromechanism in the MQ Ni–P amorphous alloys has been discussed [39]. The steep increase in the microhardness curve could be connected with the formation of Ni3P particles in addition to Ni particles. The wellinvestigated phenomenon of precipitation hardening of monocrystalline materials is characterized with linear dependence on the reciprocal radius of precipitates and
366
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
with much weaker dependence on their volume fraction. (Relatively uniform amorphous matrix could be considered as closer to monocrystalline structure, than to the polycrystalline one with a strong influence of the grain boundaries.) We did not obtain information about the mechanism of the precipitation hardening in amorphous matrix, nevertheless a similar effect was observed in amorphous alloys [40]. When the bulk crystallization takes place, the volume fraction of the crystalline Ni3(P,Ni) phase increases more than three times faster than the volume fraction of Ni. The dependence of the microhardness of the Ni –P alloys on the average particle size has a maximum [41]. The hardness of nanocrystalline Ni80P20 alloy increases from 10 to 12 GPa when the average grain size grows from 10 to 100 nm [2]. Above 70– 100 nm the microhardness drops according to the Hall–Petch equation [42]. Therefore the highest microhardness is achieved in a completely crystalline sample with average grain size about 100 nm. The peak value of 12 GPa is obtained for our hypoeutectic sample after heating to 590 K. Coming back to the magnetothermal curve it looks like a good correlation exists between DSC peaks and magnetothermal curves maxima. As it is well known, the DSC peak reflects the bulk crystallization. The problem is that the DSC peaks are shifting to higher values, when the rate of isochronal heating is increased, but magnetothermal maxima around 600 K remain at the same place [32]. One should remember, that Ni Curie temperature (627 K) is crossed while measuring both the magnetization curves: during heating and during cooling. The strong para-process in the vicinity of the Curie temperature, as well as the abrupt change of the magnetization may considerably influence the h(T) and dh(T)/dT curves. Thus, both the height and position of the most prominent peak (at 608 K in Fig. 3) in dh(T)/dT curve could be affected. We should also note that beyond the Curie temperature the estimation of the Ni precipitation and dh(T)/dT is correct, only if the paramagnetic susceptibility of Ni is several times greater than that of Ni3P. There are indications that this is true [32]. The last peak at 650– 670 K is in the paramagnetic range. However, it is not present in the MQ samples [32]. Therefore, we may consider it as a consequence of some inhomogeneity or peculiarity of the ED samples. Thus the utilization of the magnetothermal method in combination with DSC and microhardness measurements reveals new details of the magnetic phase precipitation.
4.2. The addition of a third element The addition of a third element can affect the magnetization versus temperature curve in two ways. The first
is by delaying the precipitation process. This will result in a shift of the peak to higher temperatures, when compared with the peak of an alloy with the same content of P, but without a third element. The second way is by affecting the phase to be precipitated and particularly its magnetic properties, rather than the kinetics of the process. This case for Ni–Me –P with given contents will result in changing the peak height in the magnetization versus temperature curve. Both effects are present for our samples. The third element can either be dissolved in the fcc Ni phase yielding Ni(Me) solid solution, or in the Ni3P phase or may form compounds with P. To the XRD sensitivity Cu compounds with P (such as Cu3P) were not found in Ni–Cu –P system with up to 50% Cu [21]. No data are available for the appearance of Sn–P or Ni –Sn compounds in the Ni–Sn –P system at low content of Sn. None of these compounds which might be formed during sample crystallization is FM. Thus, the only phase with strong magnetic properties that may be present in the sample, is the Ni or Ni(Me,P) solid solution. The two metals, Cu and Sn, are related differently to Ni. Copper is unlimitedly soluble in Ni, while Sn has only very limited solubility in Ni [43]. However, tin has four times stronger effect on the magnetic moment of Ni, than Cu has.
4.2.1. Addition of Sn As a whole, the magnetization versus temperature curve of the Ni–Sn–P alloys is similar to that of the Ni –P alloys (Fig. 4). If one accounts for the position and height of the peaks, then both effects of the third element discussed above are seen: the retardation of the transformation and the reduction of the magnetization due to the alloying of the precipitated Ni particles with Sn. Let us consider the MQ Ni–Sn –P alloys in Fig. 4. The precipitation process in the 0.40 at.% Sn sample begins at almost the same temperature like in the sample with 19 at.% P, not containing Sn. However, the speed of magnetization increase is not as high as for the sample without Sn and the maximum magnetization achieved is lower. The MQ sample with 0.90 at.% Sn has the highest temperature of steepest magnetization increase and further diminution in the height of the magnetization maximum. Keeping in mind that the sample without Sn contains 19 at.% P, one should suppose that the sample with 18 at.% would have even lower temperature of steepest magnetization increase. Thus, in the MQ samples, the addition of Sn retards the Ni precipitation process and therefore the crystallization of the sample. The last statements are obviously true also for the ED samples. However, the increase of both Sn content from 0.2 to 0.7 at.% and P content from 19 to 20 at.%, does not lead to an increase of the
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
367
Ni-P Fig. 6. Relative magnetization h(T) =| Ni-Me-P heating / | cooling curves of amorphous alloys heated at a rate of 3 K/min.
temperature of steepest magnetization increase. This may be due to the opposite effects of the increase of both phosphorus (see Figs. 2 and 4) and tin. The retardation of the crystallization processes due to the introduction of Sn in Ni– P observed in this work using magnetic measurements is in correlation with the earlier DSC studies [18]. A shift of the DSC peak of the ED Ni –Sn –P alloys to the higher temperature was found in comparison with its position in the case of the ED Ni – P alloy. For the ternary Ni– Me – P alloys the precipitating FM phase is not pure Ni. Therefore, the saturation magnetization varies not only with the temperature and phase quantity, but also with the content of the third element in the precipitating phase. This is true for both heating and cooling the sample. Samples with different content of third element have differing cooling curves. This renders their h(T) ratios incomparable. Thus, it is more beneficial to divide by magnetization during heat-
ing to the temperature dependence of the magnetization of pure Ni, rather than by the magnetization of the sample during cooling. Before the division, the Ni magnetization versus temperature curve should be rescaled to the quantity of pure Ni that would precipitate after complete crystallization as if the alloy was Ni–P. Then h(T) will be the ratio between the magnetic moment of the precipitating FM phase and the magnetic moment the sample would have after complete crystallization if pure Ni from Ni–P were precipitating. As can be seen from Fig. 6, the h(T) curve of the Ni –Sn–P alloys is similar to those of the Ni–P alloys, but does not tend to 1 at higher temperatures. This is due to the presence of Sn in the precipitating Ni phase. Moreover, there is a decrease above 670 K for samples with higher Sn content. This manifests in the negative values of dh(T)/dT curves in Fig. 7 and indicates that Sn incorporates in the particle to a larger extent at higher temperatures.
Ni-P Fig. 7. Differentiated h(T):dh(T) / dT= d(| Ni-Me-P heating / | cooling) / dT curves of amorphous alloys heated at a rate of 3 K/min.
368
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
As a whole, the dh(T)/dT curves of the Ni–Sn –P alloys are also similar to that of the Ni– P alloy. They appear at the same position and have similar shape, but the distinct peaks in the Ni– P alloy seem to be ‘smoothened’ in the Ni–Sn – P alloys. Thus, the addition of Sn retards the precipitation of magnetic phase, which is now Ni(Sn) instead of Ni. The lower maxima are evidence for lower rate of magnetization increase. It is also clear from Fig. 7 that the precipitating FM phase contains Sn in a quantity proportional to the Sn content in the sample. This phase is most probably a supersaturated solution of Sn in Ni, and the content of Sn in this phase varies in the course of the phase transformation.
4.2.2. The addition of Cu The addition of Cu has an even more dramatic effect on the magnetization versus temperature curve, i.e. the magnetization slightly changes from the level of the paramagnetic amorphous phase. However, the small increase in the range 570– 580 K indicates a common behavior with a high P/Sn samples, i.e. Cu retards the crystallization no more than Sn does, as was already shown by DSC measurements [18]. Copper incorporates into the precipitating Ni particles to a quantity that is at least four times larger. Table 1 shows that this is true for the fully crystallized samples. The reason for this is the higher solubility of Cu in Ni. Since there is very small change in the magnetization of the Ni –Cu – P alloys during heating the h(T) and dh(T)/dT curves were not calculated. The Ni– Cu –P alloys retained lowest magnetization also after complete crystallization and cooling to room temperature, so that the copper is a better additive if strong FM properties have to be reduced. 4.3. The magnetization during cooling From the Curie temperatures estimated through the magnetization during cooling after the annealing step at 800 K the content of Sn and Cu in the precipitated Ni particles was evaluated and is presented in Table 1. The solubility of Sn in Ni at the annealing temperature, 800 K, is about 1 at.%. For three of four samples the content of Sn in the Ni particles exceeds the equilibrium solubility limit. Thus, half an hour at 800 K was not enough for the alloys to reach the equilibrium state. Therefore the estimated Curie temperatures and the magnetization during cooling can be used to draw some conclusion about the crystallization process. The incorporation of Sn in the precipitating Ni particles seems to be proportional to the Sn content in the entire sample. Also it does not depend on the type of sample preparation: MQ or ED. The same cannot be said for the shape of the magnetization versus temperature curve during cooling. The shape of the magnetiza-
tion curve of the MQ samples is very close to that of the electroless Ni– P samples and pure Ni while the shape of the electroless Ni–Sn– P samples is slightly different. This makes us think that the addition of Sn in the ED Ni –P samples affects the kinetics of Ni particles precipitation by a way different from that of the MQ samples, which results in different final structure.
5. Conclusion The crystallization of amorphous Ni–P, Ni –Sn–P and Ni– Cu –P alloys is investigated by a modification of a magnetic phase analysis method called magnetothermal method. It consists of three steps. The first step is the measurement of magnetization during heating. Then annealing to complete crystallization of the sample is conducted and the magnetization during cooling is measured. Then the ratio between two magnetization values (during the heating and cooling) is constructed as dependent on temperature, h(T), which is the second step. The third step is the calculation of the derivative: dh(T)/dT. The last operation is making the magnetothermal curve directly comparable to DSC curves, as far as peak positions are concerned. The magnetothermal method is much more sensitive to the precipitation of Ni or Ni-based phases during isochronal heating than DSC and this is demonstrated at a heating rate of 3 K/min. In this way, the details of the precipitation processes could be investigated. Particularly, it was shown that in the hypoeutectic alloys besides the expected primary Ni precipitation, there is an additional secondary Ni precipitation. This secondary precipitation could be noticed to a smaller extent in eutectic and even in the hypereutectic alloys with phosphorus content close to the eutectic composition. The microhardness measurement during isochronal heating also reveals the Ni precipitation in amorphous Ni –P matrix. The effect of this precipitation on microhardness, as well as estimation based on magnetic measurements suggests that the precipitated particles are of nanometer size. The addition of a third element to the Ni–P amorphous alloy can affect its thermomagnetic behavior in two ways. First, by delaying the strong magnetic phase precipitation process. Second, by affecting the magnetic properties of the precipitating phase due to alloying. The addition of Sn or Cu to the Ni–P alloy leads to the incorporation of the third element into the precipitating Ni particles. Since both Sn and Cu decrease the spontaneous magnetization of Ni, they diminish the magnetization of the Ni–Me –P alloy. Simultaneously they retard the precipitation of the Ni(Me) phase. Though Sn could have a stronger effect on the Ni spontaneous magnetization, it decreases the magnetization of the
D. Tache6 et al. / Electrochimica Acta 47 (2001) 359–369
Ni –P alloy to a lesser extent than Cu does, due to the lesser solubility of Sn in Ni. However, the addition of Sn shifts the crystallization process to higher temperatures. Thus, the addition of Cu or Sn reduces the appearance of phase with strong magnetic properties. Obviously, the proposed magnetothermal method in combination with DSC and microhardness measurements could give valuable information for the kinetics of crystallization of multicomponent amorphous alloys containing a strong magnetic phase. In this way, the use of magnetic measurements becomes much more helpful in investigating the formation of magnetic nanoparticles in an amorphous para- or diamagnetic matrix, even though the interpretation of magnetization data is not always easy and straightforward.
[12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24]
Acknowledgements
[25]
The authors are indebted to Professor Dr Sc. I. Avramov from the Institute of Physical Chemistry of Bulgarian Academy of Sciences for the DSC measurements, the Copernicus program for the DSC equipment and to Mr Ralev from DZU Plc. in Stara Zagora for access to the VSM.
[26] [27] [28] [29] [30] [31] [32] [33] [34] [35]
References [1] Ts. Hentschel, D. Isheim, R. Kirchheim, F. Mueller, H. Kreye, Acta Mater. 48 (2000) 933. [2] K. Lu, Mater. Sci. Eng. R16 (1996) 161. [3] R.N. Duncan, Paper presented at COATINGS’99, Conference Proceedings, Dallas, 21 –23 September, 1999, p. 265. [4] J. Hajdu, Trans. Inst. Met. Finishing 75 (1) (1997) B7. [5] D.R. Frear, J. Met. 51 (3) (1999) 22. [6] U. Pittermann, S. Ripper, Phys. Stat. Sol. (a) 93 (1986) 131. [7] U. Pittermann, S. Ripper, in: S. Steeb, H. Warlimont (Eds.), Rapidly Quenched Metals, Elsevier, Amsterdam, 1985, p. 375. [8] U. Pittermann, S. Ripper, Z. Metall. 74 (1983) 783. [9] K.-H. Hur, J.-H. Jeong, D.N. Lee, J. Mater. Sci. 25 (1990) 2573. [10] K.-H. Hur, J.-H. Jeong, D.N. Lee, in: L. Huang (Ed.), Thin Films and Beam-Solid Inter, Elsevier, Amsterdam, 1991, p. 247. [11] E. Wachtel, I. Bakonyi, J. Bahle, N. Willmann, A. Burgstaller, W. Socher, J. Voitla¨ nder, A. Lovas, H.H. Libermann, Mater. Sci. Eng. A 133 (1991) 196.
.
[36] [37] [38] [39] [40] [41]
[42]
[43]
369
K. Lu, J.T. Wang, J. Non-Cryst. Solids 117/118 (1990) 716. K. Lu, J.T. Wang, Mater. Sci. Eng. A 133 (1991) 500. K. Lu, J.T. Wang, J. Cryst. Growth 94 (1989) 448. K. Masui, M. Tachihara, T. Yamada, T. Tsujimoto, J. Jpn. Inst. Met. 44 (1980) 124. K. Masui, S. Maruno, T. Yamada, J. Jpn. Inst. Met. 41 (1977) 1130. K. Masui, T. Yamada, Y. Hisamatsu, J. Met. Fin. Soc. Jpn. 31 (1980) 667. N. Krasteva, S. Armyanov, J. Georgieva, N. Avramova, V. Fotty, J. Electron. Mater. 24 (1995) 941. N. Krasteva, V. Fotty, S. Armyanov, J. Electochem. Soc. 141 (1994) 2864. S. Armyanov, N. Krasteva, V. Fotty, Proc. Int. Symp. GALVANO’91, Varna, 30 September – 2 October 1991, p. 240. K.-H. Hur, J.-H. Jeong, D.N. Lee, J. Mater. Sci. 26 (1991) 2037. X. Haowen, Zh. Bangwei, Y. Qiaoqin, Trans. IMF 77 (1999) 99. B.A. Apaev, Fazoviy Magnitniy Analiz Splavov, Metallurgia, Moskva, 1976 (in Russian). A. Berkowitz, E. Kneller (Eds.), Magnetism and Metallurgy, vols. 1 – 2, Academic Press, New York, 1969. E. Wachtel, N. Willmann, B. Predel, J. Magn. Magn. Mater. 59 (1986) 115. K. Lu, R. Lu¨ ck, B. Predel, J. Alloys Comp. 201 (1993) 229. R. Lu¨ ck, J. Mater. Sci. Technol. 13 (1997) 260. R. Lu¨ ck, K. Lu, W. Frantz, Scr. Metall. Mater. 28 (1993) 1071. A. Aronin, G. Abrosimova, I. Zver’kova, D. Lang, R. Lu¨ ck, J. Non-Cryst. Solids 208 (1996) 139. H.G. Zolla, F. Spaepen, Mater. Sci. Eng. A 204 (1995) 71. H.G. Zolla, F. Spaepen, IEEE Trans. Magn. 31 (1995) 3814. D. Tachev, D. Iorgov, S. Armyanov, J. Non-Cryst. Solids 270 (2000) 66. A. Amamou, J. Durand, Commun. Phys. 1 (1976) 191. J.S. Kouvel, J.B. Comly, Phys. Rev. Lett. 24 (1970) 598. I. Baconyi, A. Burgstaller, W. Socher, J. Voitla¨ nder, E. TothKadar, A. Lovas, H. Ebert, E. Wachtel, N. Willmann, H.H. Liebermann, Phys. Rev. B 47 (1993) 14961. A. Aharoni, IEEE Trans. Magn. 29 (1993) 2596. W.F. Brown, Ann. N. Y. Acad. Sci. 147 (1969) 463. E.C. Stoner, E.P. Wolhfarth, Phil. Trans. R. Soc. A 240 (1948) 599. K. Lu, J.T. Wang, Mater. Sci. Eng. A 133 (1991) 504. H. Kimura, M. Masumoto, in: F.E. Luborsky (Ed.), Amorphous Metallic Alloys, Butterworths, London, 1983 (chap. 11). R. Rofagha, U. Erb, in: H.D. Merchant (Ed.), Defect Structure, Morphology and Properties of Deposits, The Minerals, Metals and Materials Society, 1995, p. 245. Y. Tobiyama, B. Derby, J.B. Pethica, in: H.D. Merchant (Ed.), Defect Structure, Morphology and Properties of Deposits, The Minerals, Metals and Materials Society, 1995, p. 335. T.B. Massalski (Editor-in-Chief), Binary Alloy Phase Diagrams, vols. 1 – 2, ASM, Ohio, 1986.