ࡱ> 5@bjbj22 8XXm\\\\\\\p804p07$"4444444$8Rn:4\###4\\6(((#@\\4(#4(>(8)r2T\\c3 `#2470072zP;&,P;c3pp\\\\P;\c3`C r( \!44ppT("ppTTEMEPERATURE DEPENDENCE OF MICROSTRUCTURE OF (1-x)Al-xZn ALLOYS, x = 0.44, 0.48, 0.54 AND 0.62 }ELJKO SKOKO and STANKO POPOVI Physics Department, Faculty of Science, University of Zagreb, Bijeni ka cesta 32, HR-10002 Zagreb, POBox 331, Croatia E-mail address:  HYPERLINK "mailto:zskoko@phy.hr" zskoko@phy.hr,  HYPERLINK "mailto:spopovic@phy.hr" spopovic@phy.hr Dedicated to the memory of Professor Zvonko Ogorelec The change of microstructure with temperature of the title alloys has been studied in situ by X-ray powder diffraction. It has been found that the temperature dependence of microstructure of the alloys, rapidly quenched from the solid-solution temperature, Tss, to room temperature, RT, is quite different from that of the alloys slowly cooled from Tss to RT. The area between two curves showing that dependence for the given phase during the first heating from RT to Tss and first cooling from Tss to RT is much smaller for the slowly cooled alloys than for the rapidly quenched alloys. That area slightly increases with the increase of the Zn content in the alloys. The temperature dependence of microstructure of the alloys during the second heating from RT to Tss and second cooling from Tss to RT differs little from that during the first cooling from Tss to RT. The ideal equilibrium state cannot be reached either by slow cooling of the alloys from Tss to RT, or by a prolonged ageing at RT of the rapidly quenched alloys. The observed sequence of phase transitions in alloys during heating from RT to Tss is different from that which could be expected according to the phase diagram of the system Al-Zn accepted in the literature. During cooling from Tss to RT, a temperature hysteresis is observed in reversal phase transitions. PACS numbers: 61.50.-f, 64.70, 64.75.-p, 64.75.+g, 65.70.+y UDC 548.73 Keywords: Al-Zn alloys, microstructure, phase transition, solid solution, X-ray powder diffraction Introduction The system Al-Zn is very suitable for studying phase transitions and microstructure in dependence on composition and/or temperature. Zinc atoms do not form intermetallic phases with aluminium atoms due to a weak mutual interaction. The atomic radius of the Al atoms is 0.143 nm, while the atomic radius of Zn atoms amounts 0.134 nm. This difference has a great influence on the microstructure of Al-Zn alloys. An Al-Zn alloy can gradually (asymptotically) approach the equilibrium state after a prolonged ageing at, say, room temperature (RT). This process can be accelerated at an elevated temperature, say, several tens Ks above RT. In such a state the alloy contains two phases: -phase (fcc, the matrix, M) having H" 99 at% Al, and (Zn)-phase (hexagonal, the precipitates) having H" 99.5 at% Zn. One can denote the (-phase in equilibrium with the phase (Zn) as (M/) The solubility of Zn in Al increases with temperature and reaches about 67 at% at 655 K [1]. The Al-Zn alloy rapidly quenched from the solid-solution temperature, Tss, to RT is supersaturated and its decomposition takes place immediately after quenching. The sequence of precipitates appearing during a prolonged ageing is as follows: spherical Guinier-Preston zones (GPZ, fcc, coherent with M, having the diameter up to 3 to 4 nm, containing about 70 at% Zn) ellipsoidal GPZ (fcc, coherent with M) rhombohedrally distorted  EMBED Equation.DSMT4 -phase (partially coherent with M, having particles of about ten nanometers in size)  metastable -phase (fcc, partially coherent with M)  (Zn)-phase (incoherent with M, the final equilibrium precipitates having micrometer sizes). This sequence of phase transitions from one to the next type of precipitates depends on the initial composition of the alloy and on the thermal treatment [1]. The equilibrium phase diagram of the system Al-Zn has been defined on the basis of a number of papers and accepted in the literature [1, 2]. In spite of many published data on the Al-Zn system, a systematic study of the system by X-ray powder diffraction (XRD), lacking in the literature, has been undertaken. A series of Al-Zn alloys having the Zn fraction x = 0.045, 0.08, 0.15, 0.20, 0.24, 0.26, 0.35, 0.38, 0.40, 0.44, 0.48, 0.54 and 0.62, prepared from components of purity 4N, were subjected to different thermal treatments. The alloys of different composition, quenched rapidly from Tss to RT, were aged at RT or at elevated temperature, and the precipitation processes, that is, the decomposition of the supersaturated solid solutions, were followed during ageing. Also, the alloys of different composition, aged after quenching for different time intervals, were subjected to a gradual change of temperature, from RT to Tss and back to RT, and their microstructure was followed in situ by XRD [3-6]. During the decomposition of the supersaturated alloy, a gradual transition of the solid solution, ss, to the equilibrium phase, (M/), takes place. That is clearly manifested in X-ray diffraction patterns because of different unit-cell parameters of the two phases, due to different content of Zn [3]. As various precipitates, P, grow inside the matrix, M, they cause strains in the crystal lattice of the matrix. X-ray diffraction lines of the matrix, (M/P), are broader than those of pure Al [3]. Also, for the alloys approaching the equilibrium state, diffraction lines of the (M/)-phase are little broader than the ones of pure Al, because of strains in the matrix crystal lattice around the (Zn) precipitates. On the other hand, diffraction lines of (Zn) precipitates are rather sharp almost from the beginning of their appearance, indicating an unstrained crystal lattice [3]. GPZ are formed after quenching in alloys having x d" 0.48. In the alloys with a higher Zn fraction, the decomposition process is rather fast and the sequence of precipitates is shortened [3]. The decomposition rate of the supersaturated alloy depends on the quenched-in vacancies (i.e. on the quenching rate), on the initial Zn fraction and the ageing temperature. It is the ageing temperature, which strongly influences the diffusion rate of the Zn atoms inside the host Al crystal lattice (the matrix) [1, 3]. These XRD studies have shown that accurate measurements of positions and profiles of diffraction lines yield useful information, e.g. on the zinc content in the matrix, M, for different precipitates, P, on the strains occurring at the M/P interface, on the unit-cell parameter of the intermediate phase, , and of the solid solution, ss. Also, the recent investigations have shown that a correction of the phase diagram is necessary. The present work is focused on the XRD study of the temperature dependence of microstructure of four Al-Zn alloys, with a high Zn fraction, x = 0.44, 0.48, 0.54, 0.62. The results obtained in this work are discussed in connection with two previous papers [5, 6]. That is, the present work is a review of a recent XRD study of the Al-Zn alloys, with the aim to check the validity of the phase diagram of the Al-Zn system in its middle part, where the fractions of Al and Zn are comparable. Experimental The powder samples for XRD were prepared by filing the bulk alloys produced from elements of purity 4N. The samples were annealed up to 2 hours in the region of the solid-solution temperature, Tss (up to 770 K), and quenched in water at RT. The samples were wrapped in a thin Al foil, perforated with hundreds of small holes in order to increase the quenching rate, which was estimated to 105 K/s. The quenching technique was described in detail in previous papers, e.g. Ref [4]. The as-quenched alloys, having x = 0.54 and 0.62, being in a supersaturated state, were subjected to a prolonged ageing at RT. Having approached the equilibrium state, the alloys were studied by XRD. These samples were denoted as water quenched and prolonged aged, in which Zn is dominant, WQZP. The alloys having x = 0.44 and 0.48 were also rapidly quenched from Tss in water at RT and then aged at RT for one week (the alloy with x = 0.48) or for two weeks (the alloy with x = 0.44), and for 11 months (both alloys, with x = 0.44 and 0.48). The samples aged at RT for one/two week/s were denoted as water quenched samples, in which Al is dominant, WQA. The samples aged at RT for 11 months were denoted as water quenched and prolonged aged samples, in which Al is dominant WQAP. The same alloys, with x = 0.44 and 0.48, were also slowly cooled from Tss to RT over 5 days, and then aged at RT for one week (the alloy with x = 0.48) or for two weeks (the alloy with x = 0.44). These alloys were denoted as slowly cooled samples in which Al is dominant, SCA. All samples, being quenched/slowly cooled from Tss to RT and then aged at RT, tended to the equilibrium state. One may suppose that, after ageing at RT, the SCA samples were closer to the equilibrium state than the WQA samples. We wanted to find out whether the quenching from Tss to RT and prolonged ageing at RT for 11 months produced a similar effect as slow cooling from Tss to RT and short ageing at RT, that is, which samples, WQAP or SCA, were closer to the equilibrium state. It was also of interest to compare samples in which Zn is dominant, WQZP, with those in which Al is dominant, WQA and WQAP. After ageing at RT, the samples were studied by XRD using a Philips diffractometer, having a high-temperature attachment (proportional counter, graphite monochromator, radiation CuK). The samples were heated (by a platinum strip) from RT to Tss and then cooled to RT, at a rate of 2 to 3 K/min. The heating/cooling of samples was stopped at a series of selected temperatures (20 to 30) for 15 minutes in order to scan prominent diffraction line profiles. In some cases two heating and cooling cycles, between RT and Tss, were performed with the same specimen (e.g. for the alloys with x = 0.54, 0.62 (WQZP) and x = 0.48 (SCA)). Some samples were exposed to air (105 Pa) and to a low air pressure (10-3 Pa), but no effect of oxidation was observed by XRD. Two or three experiments were performed with each alloy and reproducible results were obtained. Appropriate precautions were undertaken in order to minimize systematic aberrations, which could influence the observed values of the unit-cell parameters of the phases , , ss, and (Zn) [7, 8]. Results and discussion In line with the phase diagram, accepted in the literature, one could expect the following phase transitions: for the alloys with x = 0.44, 0.48 and 0.54,  EMBED Equation.DSMT4  for the alloys with x = 0.62,  EMBED Equation.DSMT4  However, the following sequence of phase transitions has been observed for the studied alloys with x = 0.44 and 0.48 (in which Al is dominant) (Ref. [6] and the present work):  EMBED Equation.DSMT4  On the other hand, for the studied alloys having x = 0.54 and 0.62 (in which Zn is dominant), the following phase transitions have been observed [4]:  EMBED Equation.DSMT4  These results indicate that a correction in the phase diagram of the system Al-Zn is necessary. Our previous and present studies have shown that the unit-cell parameter of the ((M/()-phase, in equilibrium with the (Zn)-phase approached after a prolonged ageing of the quenched alloy, does not depend on the initial composition, amounting 0.40469(6) nm at RT. On the other hand, the unit-cell parameter of the (M/)-phase of the alloy approaching the equilibrium state by slow cooling from Tss to RT and by ageing at RT, is 0.40445(10) nm, regardless of the initial composition (Ref. [3] and references therein, Refs. [4, 6], the present work). That difference between the two values of the unit-cell parameter is due to different content of Zn retained in the matrix, as a result of different thermal treatments. For comparison, the unit-cell parameter of pure Al amounts 0.40494(3) nm [3]. From Refs. [3, 4, 6] and the present work] it follows that the unit-cell parameters of the (Zn)-phase, in equilibrium with the (M/)-phase, are close to those of pure Zn at RT, namely, a = 0.2665(2), c = 0.4947(3) nm (space group P63/mmc). The above cited values have been used to calibrate the angular scale in high-temperature XRD measurements in the previous works [4, 6] and in the present work. The figures given in the present work correspond to the alloys, quenched from Tss to RT and aged at RT for 11 months (samples WQAP). Figure 1 shows prominent diffraction line profiles (scanned in situ) at medium Bragg angles, of the alloy with x = 0.44 at selected temperatures, including the heating run and the cooling run. Figure 2 shows the temperature dependence of the unit-cell parameter a of the phases ,  and ss of the alloy with x = 0.44, during the heating and cooling runs. The direction of temperature change is shown by arrows, and vertical bars indicate the estimated standard deviation (e.s.d.) in the derived parameter values. The same features are also given in the following figures. The temperature dependence of the unit-cell parameters a() and c() of the (Zn)-phase is shown in Figs. 3 and 4, respectively, for the alloy with x = 0.44. Figure 5 shows the temperature dependence of an interplanar spacing of the (Zn)-phase, namely d103 (which is of a particular interest, including both the temperature behaviour of the a and c axes) in the alloy with x = 0.44. Finally, the temperature change of diffraction line intensities of diffraction lines 110 and 103 for the alloy with x = 0.48 is shown in Fig. 6. Similar figures showing the temperature dependence of microstructure of the alloys in the Zn-rich region (water-quenched from Tss to RT and prolonged aged at RT, WQZP) and in the Al-rich region (water quenched/slowly cooled from Tss to RT and aged for one/two week/s at RT, WQA, SCA) were given in the previous papers [4, 6]. General features of the temperature dependence of XRD patterns, and therefore of their microstructure, of the studied alloys are as follows. By comparison of the water quenched alloys (WQZP, WQA, WQAP) with the slowly cooled alloys (SCA) one can conclude that the temperature dependence of their microstructure is rather different, e.g. Ref. [6], Figs. 2, 4, 6 vs. Figs. 3, 5, 7. The two curves, which show this dependence during the first heating from RT to Tss and cooling from Tss to RT, are closer to each other for the slowly cooled samples than for the water quenched samples. That is, the area between these two curves is smaller for the slowly cooled samples than for the water quenched samples. That area slightly increases with the increase of the Zn content, x, in the alloys. For the water quenched samples, the curves corresponding to the second heating from RT to Tss and second cooling from Tss to RT are close to the one of the first cooling from Tss to RT; these curves can be approximated by a linear function, e.g. Ref. [4], Figs. 2, 3, 5; Ref. [6], Figs. 4, 6; the present work, Figs. 3-5. On the contrary, the curve corresponding to the first heating from RT to Tss is quite different, deviating much from the linear function. Generally, at a given temperature, the difference between the values of a microstructural parameter, found in the heating run and in the cooling run, is smaller for the slowly cooled samples than for the water quenched samples, this difference slightly decreases with the prolonged ageing of one and the same alloy (e.g. samples WQAP vs. samples WQA). As the temperature of the alloy increases, a decrease of the diffraction line intensities of both (M/)- and (Zn)-phases takes place, due to enhanced thermal vibrations of the atoms. A gradual shift of the diffraction lines toward smaller Bragg angles is observed, caused by thermal expansion (Fig. 1 in Refs. [4, 6], the present work). Thermal expansion of the (Zn)-phase is anisotropic. It can be followed, e.g. from the temperature dependence of the angular separation of adjacent diffraction lines, 110 and 103 (Fig. 1 in Refs. [4, 6], the present work). Thermal expansion along the c-axis is several times bigger than that along the a-axis of the (Zn)-phase (Table 1 in Ref. [4]). An indication of the thermal expansion anisotropy may also be a decrease of the saddle intensity between the adjacent diffraction lines 110 and 103 as the temperature increases (Fig. 1 in Refs. [4, 6], the present work). For the studied alloys (Refs. [4, 6], the present work) one can also observe a change of the shape of the (Zn)-precipitates with the increase of the temperature: the broadening of diffraction lines with l `" 0 (e.g. 002, 103) increases in relation to broadening of diffraction lines with l = 0 (e.g. 100, 110). (Zn)-precipitates become more and more flat, i.e. their size along the c-axis decreases as the temperature increases. Zn atoms, which leave the (Zn)-precipitates, are dominantly those which form the lattice planes {001}. This effect was not observed for the alloys with smaller Zn content (x d" 0.40 Ref. [3]). It may be supposed that the concentration of vacancies in the (Zn)-phase increases with temperature; this affects the temperature dependence of its a- and c-unit cell parameters for the water-quenched alloys (Ref. [4] (WQZP), Figs. 3-5; Ref. [6] (WQA), Figs. 4, 6; the present work (WQAP), Figs. 3-5). For the water quenched and slowly cooled alloys the unit-cell parameter, a, of the (M/)-phase changes rather linearly during the first heating run up to H" 500 K (Ref. [4], Fig. 2; Ref. [6], Figs. 2, 3; the present work, Fig. 2). At a higher temperature, a partial dissolution of Zn from (Zn)-phase into the (M/)-phase is observed. This process compensates and even reverses the change of the unit-cell parameter a of the (M/)-phase due to thermal expansion, as the Zn atoms are smaller than the Al atoms (Ref. [4], Fig. 2, Ref. [6], Figs. 2, 3; the present work, Fig. 2). Above H" 550 K a new phase,  (fcc), appears in accordance with the phase diagram [1, 2]. The unit-cell parameter of the -phase, being rich with Zn, is smaller (0.8 to 0.9 % at 560 K) than the one of the -phase. Therefore, diffraction lines of the -phase are at the high-angle side of diffraction lines (having the same Miller indices) of the -phase. In line with the phase diagram, for the alloys with x = 0.44, 0.48, and 0.54, the -phase should be in coexistence with the -phase above H" 550 K, i.e. the (Zn)-phase should completely transform into the -phase above that temperature. On the other hand, for the alloy with x = 0.62, the phases  and (Zn) should be present above H" 550 K. However, our previous [4, 6] and present results indicate that only a partial transition of the (Zn)-phase, and probably a partial transition of the (M/)-phase, into the  (fcc)-phase takes place above H" 550 K. It follows that the -phase is in coexistence with both - and (Zn)-phases, and it is now denoted as (M/, ). As the temperature further increases, the composition of the phases changes, the content of Zn in the (M/, )-phase, and probably in the -phase, increases. These Zn atoms come from the (Zn)-phase, in which the concentration of vacancies increases; t`bjl   * + , 9 : < = a b c r s u  G vjjh !hACH*mH sH h !hBCmH sH h1QhAC5mH sH #jh !hACUmH sH h !hAC0JmH sH #jh !hACUmH sH jh !hACUmH sH h !h)mH sH hj;mH sH hj;6mH sH h !hAC6mH sH h !hACmH sH &  t u /0|d&y)-,$ & Fd^a$gd, $da$gd, $da$gd,NG [  ijl (lnʾʾʾʾʾʾʪʟujhDyhDymH sH  jahDymH sH h1Qh6uCJaJmH sH h1Qh6u6CJaJmH sH h !h6umH sH h\lmH sH h !hBCH*mH sH h !hBC6mH sH h !hBCmH sH h !hACH*mH sH h !hAC6mH sH h !hACmH sH hDymH sH )nzabdJq Nʓʋth\tTth\thmH sH h !he H*mH sH h !he 6mH sH h !he mH sH h !hW6mH sH hmH sH h37dmH sH !j~h !hk_EHUmH sH jlE hk_CJUVaJjh !h6uUmH sH h !hWmH sH h !h6uH*mH sH h !h6u6mH sH h !h6umH sH hDymH sH  NUc $&&&&T'V'y)*P*++,,,2----...#.$.%.8.9.;.ƻƯƻ|l]|U|UhywmH sH h1Qh`$CJaJmH sH h1Qh`$6CJaJmH sH h !h`$mH sH hDymH sH h !hn6mH sH h !hnH*mH sH h_>mH sH h !h1`6mH sH h !hnmH sH h !h1`mH sH h !he H*mH sH h !hWmH sH h !he mH sH h !he 6mH sH !-,..$.%.031;46<<<<==>>8>V>r>"? $da$gd,$ & F 8d^`a$gd $da$gd$ & Fd^a$gd, $da$gd,;.h.p......///0$0%000121E1F1w1x1z111111122v22238393h3i3k33333%434i4j4k4m4x4}4h !hH*mH sH h !h6mH sH h !hmH sH hDymH sH hk_mH sH hmH sH h !h`$H*mH sH h !h`$H*mH sH h !h`$6mH sH hmH sH h !hZmH sH h !h`$mH sH 3}45,5Q5R5T5555556g6h666P78b8d8z8|88X9j9P:Q:S:z:|:::::::::; ;;;;;<º¯£›£‡£££{{pph !hZmH sH h !hp0H*mH sH hDymH sH h !hp0H*mH sH hmH sH h !hp06mH sH h !h`$mH sH hk_mH sH h !hp0mH sH hmH sH h !hH*mH sH h !h6mH sH h37dmH sH h !hmH sH ,<<<<<=>>>>>4>5>6>7>8>L>M>V>W>n>o>p>q>r>>>>>>農xiXMh !hZmH sH !j7h !hDyEHUmH sH jH hDyCJUVaJ!jkh !hDyEHUmH sH jH hDyCJUVaJjh !h}UmH sH h !h}mH sH h !h6mH sH h !hmH sH h1Qhp0CJaJmH sH h1Qhp06CJaJmH sH h !hp0mH sH h !hp0H*mH sH >? ?? ?#?$?;????q?r????????????JALARA˼}n]}O jah !hmH sH !jh !hDyEHUmH sH j3H hDyCJUVaJh !hBmH sH jh !hBUmH sH hk_mH sH h !h6mH sH !j h !hDyEHUmH sH j͢H hDyCJUVaJh !h}mH sH jh !h}UmH sH h37dmH sH h !hmH sH hmH sH "?#???@?????@rED@DADZDaDeDyDRErEzE{EEEFFFFF$G&G(G*G,G2G;HmH sH hmH sH h !hmH sH  jbh !hmH sH .IIIJJ@JBJJJJNKKNLLLLLL&M0M2MxMzMMMMMMMMMNJNLNRNNOO OO0OOLOPLPrPrQtQxQQ R!R#R-R8RYReRiRRúúúòúúúúúúúhY_mH sH h !H*mH sH hZmH sH h !6mH sH h !mH sH h$mH sH hDymH sH h !h !mH sH hH*mH sH h6mH sH hmH sH PmH sH hRmH sH hmH sH hjmH sH >&e6e@eBee;fmH sH h_Z`H*mH sH h_Z`6mH sH h]mH sH h)mH sH h>PmH sH hk_mH sH hHnmH sH h_Z`mH sH hmH sH h'SmH sH hv8mH sH 0 0vx|Ԣ ">VDHƦЦҦԦV LVXZ\b,Bսͽͽͽͽͽͽͽͽͽͽͽսh1Qh!4L6CJaJmH sH hmH sH h!4L6mH sH h]mH sH h!4LmH sH hRmH sH h'SmH sH hmH sH hTj6mH sH hH^mH sH hTjmH sH hTjH*mH sH :B\+ghjȮ<?$jmn*BflżδάΤ|w|s|n h.m5h] h.m6h.mUh1QmH sH h.mmH sH hRH*mH sH hR6mH sH hRmH sH h]mH sH hH^mH sH hKH*mH sH hK6mH sH hKmH sH h!4LH*mH sH h!4L6mH sH hHnmH sH h'SmH sH h!4LmH sH +n (*2B$ & F d^`a$gd, $da$gd, $da$gd, References H. Lffler, Structure and Structure Development in Al-Zn Alloys, Akademie Verlag, Berlin (1995). H. Lffler, G. Wendrock and O. Simmich, Phys. Stat. Sol. (a) 132 (1992) 339. S. Popovi and B. Gr~eta, Croat. Chem. Acta 72 (1999) 621-643 and references therein. S. Popovi, B. Gr~eta, B. Han~ek and S. Hajster, Fizika A 8 (1999) 173. S. Popovi and B. Gr~eta, Mater. Sci. Forum 321-324 (2000) 635. }. Skoko and S. Popovi, Fizika A 10 (2001) 191. S. Popovi, J. Appl. Cryst. 4 (1971) 240-241, 6 (1973) 122. S. Popovi, Cryst. Res. Technol. 20 (1985) 552. TEMPERATURNA OVISNOST MIKROSTRUKTURE SLITINA (1-x)Al xZn, x = 0.44, 0.48, 0.54 AND 0.62 Ovisnost mikrostrukture navedenih slitina o temperaturi istra~ivana je in situ pomou rentgenske difrakcije u prahu. Pokazano je da je temperaturna ovisnost mikrostrukture slitina, brzo kaljenih s temperature vrste otopine, Tss, na sobnu temperaturu, RT, bitno razli ita od one za slitine, koje su sporo hlaene s Tss na RT. Povraina izmeu dviju krivulja, koje pokazuju tu ovisnost tijekom prvog grijanja slitine od RT do Tss i prvog hlaenja od Tss do RT, mnogo je manja za sporo hlaene slitine nego za brzo kaljene slitine. Ta povraina lagano raste s udjelom Zn u slitinama. Temperaturna ovisnost mikrostrukture slitina tijekom drugog grijanja od RT do Tss i drugog hlaenja od Tss do RT malo se razlikuje od ovisnosti tijekom prvog hlaenja od Tss do RT. Idealno ravnote~no stanje ne mo~e se postii niti sporim hlaenjem slitina od Tss do RT niti dugim starenjem pri RT slitina brzo kaljenih od Tss na RT. Opa~eni niz faznih pretvorbi u slitinama tijekom grijanja od RT do Tss razlikuje se od onoga koji bi se o ekivao prema faznom dijagramu sustava Al-Zn, prihvaenog u literaturi. Tijekom hlaenja slitine od Tss do RT uo ena je temperaturna histereza za obrnute fazne pretvorbe. FIGURES Fig. 1. Prominent diffraction lines of the alloy Al-44 at% Zn at selected temperatures. The alloy was quenched from the solid-solution temperature, Tss, to RT (in water) and aged at RT for 11 months. Radiation: monochromatized (graphite) CuK(, counter: proportional. Fig. 2. The dependence of the unit-cell parameter, a, of the phases (, ( and (ss in the water-quenched alloy Al-44 at% Zn (aged at RT for 11 months) on temperature during heating from RT to TSS and cooling to RT. The arrows indicate the sense of temperature change. Vertical bars indicate the estimated standard deviation (e.s.d.) in a. Fig. 3. The dependence of the unit-cell parameter a(() of the phase ((Zn) in the water quenched alloy Al-44 at% Zn (aged at RT for 11 months) on temperature during heating from RT to Tss and cooling to RT. The arrows indicate the sense of temperature change. Vertical bars indicate e.s.d. in a((). Fig. 4. The dependence of the unit-cell parameter c(() of the phase ((Zn) in the water-quenched alloy Al-44 at% Zn (aged at RT for 11 months) on temperature during heating from RT to Tss and cooling to RT. The arrows indicate the sense of temperature change. Vertical bars indicate e.s.d. in c((). Fig. 5. The dependence of the interplanar spacing d103 of the phase (Zn) in the water-quenched alloy Al-44 at% Zn (aged at RT for eleven months) on temperature during heating from RT to Tss and cooling to RT. The arrows indicate the sense of temperature change. Vertical bars indicate e.s.d. in d103. Fig. 6. The dependence of peak intensity (arbitrary units) of diffraction lines 110 and 103 of the phase ((Zn) on temperature for the alloy Al-48 at% Zn during heating and cooling cycle. The alloy was water quenched from TSS to RT and aged at RT for 11 months. The arrows indicate the sense of temperature change. Vertical bars indicate e.s.d. in peak intensity. PAGE  PAGE 17 l02(@B^bz|~ B źźhNw:h$4H*hNw:h$46 hNw:h$4hj;6CJaJh1Qh$46CJaJh1Qh$4CJaJhj;CJaJh"yh$4h$4h.mmH sH  h?z5h?zmH sH  h?z5\h?zh]h.m8|~BD$a$gd, $da$gdr$ 7da$gd,$ & F d^`a$gd,hjn "&,PRV&*\ *<Pźźźŧźźź jah(mH sH hDvhDv6mH sH hrH*mH sH hDv6]mH sH hDvmH sH h(6]mH sH h(mH sH h"ymH sH h1Qh(h$4 hNw:hNw:hNw:h$46 hNw:h$4hNw:h$4H*3lnvxN`bd|~h]hgdr &`#$gdr$a$gd,$a$gd,PRVXdfj:>HLNlpnz:>@DFNRTû֘vv jbh(6]mH sH h6]mH sH h(H*mH sH hDv6]mH sH hDvhDv6mH sH hDvhH^6mH sH hH^mH sH hH^6]mH sH hDvmH sH h(mH sH hrH*mH sH h(6]mH sH  jah(6]mH sH .,.LNPT &*48X\&~ ~hH^6mH sH hXmH sH hhXH*mH sH hXhXmH sH h6mH sH hX6mH sH hrH*mH sH hX6]mH sH hmH sH h(mH sH h6]mH sH  jbh(6]mH sH h(6]mH sH ,jrvxJLNR&48@D^blr~LN¹°¤››’›‡››’|sgh1Qh(6mH sH hH^6mH sH hX6]mH sH h(H*]mH sH hX]mH sH h(]mH sH  jbh(6mH sH hHn6mH sH h6mH sH h(6mH sH hXhXmH sH hX6mH sH hXmH sH hXhXH*mH sH hXhX6mH sH 'NP\^`dfrtxz|h1Qh(6mH sH h]h]0JmHnHuh37d h37d0Jjh37d0JU6&P 1h:pj;. A!"#$% DyK zskoko@phy.hryK *mailto:zskoko@phy.hrDyK spopovic@phy.hryK .mailto:spopovic@phy.hrDd @hb  c $A? ?3"`?27 ׎n۰ZQ-`!  ׎n۰ZQ@ |xڍRMKQ=#D-"Jf!}dE#!t9u)$7tMt޽ @(.J)flKZ\;.a> pl(F$0Th&ذKe1P&#O.m8JtV}R,COH4&NSnL/Q5vf)7^2,9;rRG9E܇_Ǐ&S-R$kTv?ǁ'ՠ'\ 淆WHX+׵=^_uڵ7[rv6҉Z|t"|ߪ-t_>vZF08hu f)05,|SD`AwfdyV0g]ɱsvzX 4(=,i܅IQۖ}9]0 ~G\Nw枩xې${u=w_&DdDd 'b  c $A? ?3"`?2oF璵 8,2-`!oF璵 8,2@%hxڥTMLQyۢ]00)  h` HMUC5ؓ#ԃ`1l䁃&r3@Z1zPnKa;{og!~ @cO B`x8V-J~5"-j C=zAqETjJFT# F{,a[dKA*OHE/#y܂<_}G_,xIhNrrps#pU ܪ~o@tWt,9+x':$aT2ow%Wu]GCmu+a5mfd3`W]w}?gӏꙎ !MƧ=ziN;z5HRTC3JdL՘kd3Vs;&kݲor}cڄH٫h7*l =DWxOt̓œ.7 fi %(xya~IƟ7iDd Lb  c $A? ?3"`?2 y{T`s[cY;q{-`!y{T`s[cY;qXu#hxڥTMLQyVT-B46b4P7.P%1%L4UVm`O񠖘z.tw:he c`ׄXO;+^6D3Wi`Z?cb%/f]s=M\HZNx#?gd1a&EpT}@_ֈ}lWxku+&Ol8kmBtБlMKG:C`. up,1e}.&kaUuJh,10XrkX6RWSch,|++e N`p> j)`DS@οDvv*zŭ8az&qv~q65jar[z[D3〝cI9é)%0oηChvS1%?T m8Ⓝ]:d8!dTkȨVWN|c썌Dɒ )!EHJՊ<= zn|ܼ>㦟'"?`}'IFyWf}t 7KE5 '4kTmovءx!PzCY5$+Bu ~r~/[Dd b  c $A? ?3"`?2:Й Z{`k= -`!Й Z{`k``/hxڥUMLA~3Ӣ]4!" %Do\ʥ`¥V]?`/Zx1x5g ӊ1&D !Bfͷ3o!``e %@E HVKK~u4K/6#:Q#B3[^ ~ p G^uMF2c2Q&DkԶQ.O`zLR:.現O19j*[бQt h-{ c88 5&[:E9 tM]gN:ګ`Bǟti4#PFWt"lrcē'XjO?x2 Dg#'Agɮ<0ϲu|8XFRyn|2L#`,g7ЌV>?y3&dXk]9Pb#^a'p911nef߈[R "VV;OIDwL_I_bdUp2Cړ1y'8-i`'!JG%PCq5E{x hZNc}FC*%D*"V3'Ӧj*M VnDLMnRmֲB-w^{:$v)s"MzxiIP~P?`(Ӿ62Q&$|}Xʴj֚- Dd   !"#$%&'()*+,-./0123456789:;<=>?@ABCDEFGHIJKLMNOPQRSTUVWXYZ[\]^_`abcdefghijklmnopqrstuvwxy{|}~Root Entry FData z WordDocument8ObjectPool `_1167570540F``Ole CompObjiObjInfo  #$%&'(),/012345689:;<=>@ABCDEFGHIK FMathType 5.0 Equation MathType EFEquation.DSMT49q:8 DSMT5WinAllBasicCodePagesTimes New RomanSymbolCourier NewMT Extra!/ED/APG_APAPAE%B_AC_AE*_HA@AHA*_D_E_E_A   REquation Native _1222156975 F``Ole CompObj i FMathType 5.0 Equation MathType EFEquation.DSMT49q DSMT5WinAllBasicCodePagesTimes New RomanSymbolCourier NewMT Extra!/ED/APG_APAPAE%B_AC_AE*_HA@AHA*_D_E_E_A  M/(ObjInfo Equation Native  _1222156993F``Ole )+Zn()!+M/()! ss ; FMathType 5.0 Equation MathType EFEquation.DSMT49q DSMT5WinAllBasicCodePagesCompObjiObjInfoEquation Native _1222157005F``Times New RomanSymbolCourier NewMT Extra!/ED/APG_APAPAE%B_AC_AE*_HA@AHA*_D_E_E_A  M/()+Zn()!+Zn()! ss . FMathType 5.0 Equation MathTyOle CompObjiObjInfo!Equation Native "pe EFEquation.DSMT49q DSMT5WinAllBasicCodePagesTimes New RomanSymbolCourier NewMT Extra!/ED/APG_APAPAE%B_AC_AE*_HA@AHA*_D_E_E_A  M/()+Zn()!+Zn()+M/,()! ss . FMathType 5.0 Equation MathType EFEquation.DSMT49q  DSMT5WinAllBasicCodePagesTimes New RomanSymbolCourier NewMT Extra!/ED/AP_1222157107F``Ole *CompObj+iObjInfo-Equation Native .(1Tablel;SummaryInformation(7DocumentSummaryInformation8?G_APAPAE%B_AC_AE*_HA@AHA*_D_E_E_A  M/()+Zn()!+Zn()+M/,()!+Zn()! ss .Oh+'08b  c $A? ?3"`?2j7/,#UF--`!>7/,#U&<h xڥVMLA~3Vٶ`"@TJxg8+RK-z-& ςalس#,R aa)DY!Lƕ,f2.7i@5o|e0pQ!H6]JD^Zsǿ8_S8;O}\GpQtLhM\YjysnF[FΤnU f :8)hq^v :UA}:kVGH4xGGGU_n3"z6Us}b'X .TˁgK~aS+`3VfnW{7cGְ9:5,$ՅN BhNmǫ,mq`J9T Qy(Svf)͠^P86أٱ;4JNcAptq ~s@ >uُ ,@ Xd    ETEMEPERATURE DEPENDENCE OF MICROSTRUCTURE OF (1-x)Al-xZn ALLOYS, x=0REME Zeljko SkokoDEPeljelj Normal.doto eljko SkokoDEP4ljMicrosoft Word 10.0@@[@xp&@@ Y2]՜.+,D՜.+,< hp  Fizicki zavod-PMFqn7TmA ETEMEPERATURE DEPENDENCE OF MICROSTRUCTURE OF (1-x)Al-xZn ALLOYS, x=0 Title4(RZ* _PID_HLINKS MTWinEqnsA > mailto:spopovic@phy.hrXlmailto:zskoko@phy.hr   FMicrosoft Word Document MSWordDocWord.Document.89q@@@ NormalCJ_HaJmHsHtHDA@D Default Paragraph FontRi@R  Table Normal4 l4a (k@(No List6U@6 AC Hyperlink >*B*phB'B  !Comment ReferenceCJaJ<@<  ! Comment TextCJaJ@j@  !Comment Subject5\H2H  ! Balloon TextCJOJQJ^JaJ<+B< ) Endnote TextCJaJ>*Q> )Endnote ReferenceH*4 @b4 rFooter  p#.)@q. r Page Numberm _`tu/0 N,+!N$i(j((((());)Y)u)%*&*B*C*****X+u.0#6c<CELLNO_RTTUUR]^^^^^^^^^^^^^^^^^Z___E````"a#a$a%a&a'a(a)a*a+a,a-aaa+f,f-f.f/f0f1f2f3f;fRAIR[&eBlPN;>?@BCDEGHIJK^_`qsuvwx-,"?n<AFart=+9<br 4 6 )7)9)Y)q)s)&*>*@****mXX::::: !!8@0(  B S  ? ) ) ) ) *) d) ) ) T) ) )  BZm  B Zm *isiresearchsoft-com/cwywcitation8*urn:schemas-microsoft-com:office:smarttagstime9 *urn:schemas-microsoft-com:office:smarttagsplace= *urn:schemas-microsoft-com:office:smarttags PlaceType= *urn:schemas-microsoft-com:office:smarttags PlaceName8*urn:schemas-microsoft-com:office:smarttagsCity:*urn:schemas-microsoft-com:office:smarttagsStreetB*urn:schemas-microsoft-com:office:smarttagscountry-region;*urn:schemas-microsoft-com:office:smarttagsaddress 258HourMinute  `````` a!amm14 be),"#"##@%C%S&V&++c-}-u....113333==|>>>>????a@c@AABBKJMJNNOOP"PSSTT=U@U[[^^^Y_Z_____D`E```````a"a-aaa*fgjTj.mHn:Fr6uDvm>PYz}9K?z]\l"yRX}B$T%'S_>Y?H^W)Dy$4v8<n`$yw(/UVBCZ@!a!a!a!aN N N N%&()+,-./1235<>?ABCJKOPQSUWXZ\^m@,@8@ "P@*X@8t@<|@@B@F@JLN@Z^`d@h\@t@UnknownGz Times New Roman5Symbol3& z Arial5& zaTahoma"1bcbY2]7Y2]7#24dTmTm 3QH)?9DTEMEPERATURE DEPENDENCE OF MICROSTRUCTURE OF (1-x)Al-xZn ALLOYS, x=0 Zeljko Skoko }eljko Skoko    CompObjJj