Effect of heat treatment and Nb__ and H contents on the phase transformation of N18 and N36 zirconiu

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2205双相不锈钢热处理工艺

2205双相不锈钢热处理工艺2205双相不锈钢热处理工艺一、引言2205双相不锈钢是一种具有优异耐蚀性和良好机械性能的材料。

为了进一步提高其性能，热处理工艺在制造过程中起着重要的作用。

本文将详细介绍2205双相不锈钢热处理工艺。

二、热处理的目的热处理是通过控制材料的加热和冷却过程，改变其组织和性能。

针对2205双相不锈钢，热处理的目的主要有以下几点：1.提高材料的硬度和强度；2.改善材料的耐腐蚀性；3.调整材料的组织，使其具备良好的韧性。

三、热处理方法热处理方法是指将材料加热到一定温度，保持一段时间后进行冷却的过程。

针对2205双相不锈钢，常用的热处理方法有以下几种：1. 固溶处理固溶处理是将材料加热到固溶温度，使固态溶解后冷却。

固溶处理能够提高材料的硬度和强度，使奥氏体和铁素体达到均匀分布。

2. 淬火处理淬火处理是将加热到固溶温度的材料迅速冷却。

淬火处理能够提高材料的硬度和强度，但会降低材料的韧性。

3. 固溶处理 + 冷镇淬火固溶处理后，通过冷镇淬火来增加材料的硬度和强度，并保持一定的韧性。

四、热处理参数的选择热处理参数的选择对于2205双相不锈钢的性能影响重大。

以下是一些常用的热处理参数选择建议：1.加热温度：一般选择在摄氏度之间，以保证奥氏体和铁素体达到均匀分布。

2.保温时间：取决于材料的厚度和要求的组织性能，一般在30分钟至2小时之间。

3.冷却速率：冷却速率越快，材料的硬度和强度越高，但韧性会降低。

冷却介质选择时，要考虑到工艺要求和材料的成本。

五、热处理后的性能与应用经过适当的热处理后，2205双相不锈钢具备较高的硬度、强度和耐腐蚀性，适用于各种要求高强度和耐腐蚀性的场合。

例如在化工、海洋工程和石油钻采业等领域得到广泛应用。

六、结论热处理工艺对于提高2205双相不锈钢的性能至关重要。

固溶处理、淬火处理以及固溶处理加冷镇淬火等方法，能够显著提高材料的硬度、强度和耐腐蚀性。

在选择热处理参数时，需要考虑到材料的要求和工艺的可行性。

211241022_镀锌热成形钢表面颜色及氧化物形成规律

表面技术第52卷第5期镀锌热成形钢表面颜色及氧化物形成规律卢岳，张彩东，齐建军，孙力，马成，刘艳丽，熊自柳（河钢材料技术研究院，石家庄 050023）摘要：目的研究保温时间对热成形钢镀锌层颜色及氧化物组成的影响。

方法通过改变镀锌热成形22MnB5钢热处理保温时间，利用色差、辉光实验、X射线光电子能谱、粗糙度检测、扫描电子显微镜和透射电子显微镜对镀层表面及截面进行观察，利用电子探针进行元素分析，研究保温前后镀层表面氧化物形貌及镀层元素分布规律。

结果随着保温时间的增加，色差值ΔE逐渐增大。

当温度处在945 ℃时，镀层连续性受到破坏，逐渐脱落。

880 ℃加热过程后，镀层表面由排列均匀连贯的圆球状氧化物组成，连续覆盖表面，且呈聚集存在趋势，镀层表面氧化物厚度出现明显差异。

当热加工时间超过6 min后，氧化物明显增多，表面厚度起伏大，呈现出不均匀分布趋势，裂纹萌生，并逐渐加深扩散。

随着加热时间的增加，整体Zn浓度有降低的趋势。

结论镀层表面主要由ZnO、FeO、Al2O3组成，ZnO连续铺满表面，并呈现连续分布的趋势，有效避免了在高温下镀层表面Zn的挥发。

保持Zn含量在一定范围内，使得镀层具有阴极保护的作用。

关键词：热成形钢；GI镀层；色差；辉光；TEM；氧化物；粗糙度中图分类号：TG306 文献标识码：A 文章编号：1001-3660(2023)05-0208-10DOI：10.16490/ki.issn.1001-3660.2023.05.020Surface Color and Oxide Formation Rule of Galvanized Hot-formed Steel LU Yue, ZHANG Cai-dong, QI Jian-jun, SUN Li, MA Cheng, LIU Yan-li, XIONG Zi-liu(HBIS Material Technology Research Institute, Shijiazhuang 050023, China)ABSTRACT:This paper took the hot-dip galvanized hot-formed steel 22MnB5 as the research object. The morphology and element composition of the surface oxides in hot-dip galvanized hot-formed steel under different heat treatment time was compared. The mechanism of the surface oxides of the hot-dip galvanized steel with the holding time was summarized, to provide a practical reference for subsequent coating and welding production.22MnB5 galvanized sheet with a thickness of 1.4 mm was used and processed into a sample of 60 mm× 60 mm. Before the test, acetone was used to clean the oil stains and attachments on the surface of the sample. The sample and place was dried in an oven. The SX2-16-13 box-type resistance furnace was used for the heat treatment. The initial heating temperature of the heat treatment was 850 ℃. The sample was taken out and water-cooled immediately. The experiment was carried out under standard atmospheric pressure, the ambient temperature was 25 ℃ and the humidity was 10%. The sample was prepared into a metallographic sample of 10×10 mm with a inlaid cross section, and was then ground and polished. The metallographic sample was corroded with 4% nitric acid alcohol solution for about 15 s. The color difference experiment was carried out on the收稿日期：2022–03–31；修订日期：2022–08–15Received：2022-03-31；Revised：2022-08-15作者简介：卢岳（1995—），男，硕士，工程师，主要研究方向为钢铁焊接及镀层材料开发。

铁酸铋2.pdf

The Effect of Heat Treatment on Structural and Multiferroic Properties of Phase-Pure BiFeO 3M.YASIN SHAMI,1,3M.S.AWAN,2and M.ANIS-UR-REHMAN 11.—Applied Thermal Physics Laboratory,Department of Physics,COMSATS Institute of Infor-mation Technology,Park Road,Islamabad 44000,Pakistan.2.—Center for Micro and Nano Devices,Department of Physics,COMSATS Institute of Information Technology,Park Road,Islamabad 44000,Pakistan.3.—e-mail:m_yasin_shami@Phase-pure multiferroic BiFeO 3(BFO)was prepared by the coprecipitation technique using diverse precursors bismuth oxide at temperature as low as 400°C.The dependence of structural,microstructural,thermal,electrical (AC and DC),and magnetic properties on sintering temperature was system-atically investigated.Uniaxially pressed samples (Ø8mm)were sintered in air at 500°C to 800°C for 4h.X-ray diffraction analysis was used to determine the amorphous and perovskite nature of as-synthesized and calcined/sintered samples,respectively.The crystallite size of sintered powders increased from 47nm to 67nm.Scanning electron microscopy showed grain growth during sintering,which improved intergranular connectivity and decreased porosity in the samples.The ferroelectric to paraelectric Curie transition temperature (T C )of pure BFO powder was detected by differential scanning calorimetry analysis and found to be 820°C ±1°C.The samples exhibited high AC resis-tivity and dielectric constant,and low loss tangent values.The samples exhibited weak ferromagnetic behavior with an unsaturated magnetization versus magnetic ﬁeld hysteresis loop at room temperature.Ferroelectric behavior and variation in remnant polarization and coercivity were observed from polarization versus electric ﬁeld loops.Enhanced capacitance in the magnetic ﬁeld revealed the magnetoelectric effect in the samples.Key words:Multiferroics,BiFeO 3,sintering,transition temperature,PE,magnetoelectricINTRODUCTIONSingle-phase multiferroic materials having mag-netoelectric coupling between ferroelectric and fer-romagnetic orders have attracted intensive research due to their fascinating fundamental physics and mesmerizing applications.These ferroic orders hardly coexist due to three-dimensional (3D)orbital electronic structure,being essential for magnetism and to reduce the tendency for ferroelectricity.1Such systems are rare in nature,but examples can be found in perovskite oxides such as BiFeO 3(BFO),BiMnO 3,and YMnO 3.BFO is a single-phase mul-tiferroic material at and above room temperature.Itsimultaneously possesses a ferroelectric phase with Curie temperature $820°C and G-type anti-ferromagnetic order with Ne´el temperature (T N )$360°C.2It is anticipated that such materials will play a promising role in next-generation memory devices,due to the intrinsic coupling between the electric polarization and magnetic moments,which could,in principle,permit data to be written elec-trically and then read magnetically.Other fasci-nating applications can be found in spintronic devices,sensors,multiple-state memory ele-ments,electrically tunable ferromagnetic resonance devices,optical ﬁlters,satellite communications,and optoelectronic devices.3,4BFO has a distorted perovskite pseudocubic crystal structure with space group R 3c .5Two per-ovskite pseudocubes aligned through their body(Received December 27,2011;accepted May 9,2012;published online May 31,2012)Journal of ELECTRONIC MATERIALS,Vol.41,No.8,2012DOI:10.1007/s11664-012-2138-y Ó2012TMS2216diagonal along[111]direction form a rhombohedral, which combines to form a hexagonal unit cell.The oxygen octahedral is rotated clockwise and anti-clockwise around the[111]direction alternately. Superimposed on this antiferromagnetic order there is a cycloidal spin order with period of620A˚.This spin has been found to be incommensurate with the structural lattice,ultimately exhibiting weak fer-romagnetism at room temperature.6The energy of the BFO unit cell can be lowered by off-centering the bismuth ion with respect to the neighboring oxygen ions.This leads to displacement of bismuth (Bi)cation relative to Fe-O6octahedra,which causes the intrinsic electric polarization.7Practical applications of BFO in industry are hindered due to the formation of impure phases during the synthesis process and volatization of Bi at elevated sintering temperatures.Conductivity in BFO is strongly affected by leakage currents resulting from oxygen vacancies and reduction of Fe3+to Fe2+ions.8Moreover,secondary phases such as Bi2O3,Fe2O3,Bi2Fe4O9,and Bi25FeO39make it very difﬁcult to achieve superior multiferroic orders in BFO.9Synthesis of phase-pure BFO is rather difﬁcult due to the kinetics and thermodynamic properties of the Bi2O3–Fe2O3system.10Recently, we have used a diverse precursor bismuth oxide and optimum conditions to synthesize phase-pure BFO nanopowders by the coprecipitation method.11Dif-ferent researchers have implemented divergent characterizations in various ranges of sintering temperatures.Yao et al.12prepared BFO by solid-state reaction method having a second phase and investigated the effect of sintering in the limited range from815°C to845°C.Bernardo et al.13stud-ied the effect of sintering temperature on BFO dopants(W6+,Nb5+,and Ti4+),carrying out struc-tural and microstructural analyses only.Ding et al.14performed various characterizations on BFO and lanthanum-doped BFO above850°C.In this work,effects of heat treatment on phase-pure BFO samples have been studied.Amorphous powder synthesized at80°C was calcined at400°C. The dependence of structural and multiferroic properties on the sintering temperature of the samples(500°C to800°C)was investigated system-atically.Variations in the capacitance and resistiv-ity of the sintered samples as a function of temperature were also observed.EXPERIMENTAL PROCEDURES Amorphous BiFeO3powder was synthesized by the coprecipitation technique using analytical-grade iron nitrate nonahydrate Fe(NO3)3Æ9H2O and com-mercial-grade bismuth oxide Bi2O3(99.98%pure). Aqueous solution of 1.0M(molar)of sodium hydroxide(NaOH)was used as the precipitating agent.For the desired crystallite size and grain morphology,a pH value of13was maintained at30°C.15Experimental details are presented elsewhere.11As-prepared powder was calcined at 400°C±5°C for3h and uniaxially pressed (10MPa)into green pellets(Ø8mm).Later,these pellets were sintered at500°C±5°C,600°C±5°C, 700°C±5°C,and800°C±5°C for4h,being named samples S1,S2,S3,and S4,respectively.All heat treatments were carried out in air using a box furnace.Structural analysis and phase composition were determined using x-ray diffractometry(XRD;PAN-alytical,X’Pert Pro)using a Cu K a radiation source, operated at40kV and30mA.Grain size and sur-face morphology were observed using scanning electron microscopy(SEM;Hitachi SU-1500) equipped with energy-dispersive x-ray spectroscopy (EDX).Thermal analysis and phase stability of the samples were studied using differential scanning calorimetry(DSC;STA409C,Netzsch,Germany). The samples were scanned from room temperature to1000°C at heating rate of10°C/min using an empty Al2O3crucible as reference.Measurements were carried out under dynamic atmosphere of inert gas(argon).The dielectric behavior,tan(d),electri-cal AC conductivity,and capacitance in constant magneticﬁeld were measured in the frequency range from20Hz to3MHz.Capacitance data as a function of temperature(30°C to400°C)at four different frequencies were also collected.These measurements were executed on a Wayne Kerr 6440B precision component analyzer.Electrical DC resistivity of the sintered samples was measured using a Keithley6487picometer,from room tem-perature up to300°C.Ferroelectric measurements were carried out using a ferroelectric tester and Keithley8512A electrometer.Magnetic properties were performed on a vibrating-sample magnetome-ter(VSM;Quantum Design Versa Lab plus).RESULTS AND DISCUSSION Structural PropertiesPhase formation,crystallite size,and structural changes of all samples were determined from XRD analysis.Figure1showed XRD patterns of the as-synthesized,calcined,and sintered BFO sam-ples.All peaks were labeled by comparing XRD data with Joint Committee on Powder Diffraction Stan-dards(JCPDS)cards no.01-072-7678(BiFeO3),00-001-1053(Fe2O3),and00-020-0836(Bi2Fe4O9). Figure1a conﬁrms the amorphous nature of as-synthesized powder in which crystalline peaks were absent in the entire2h range.The powder crystallized into perovskite phase(BiFeO3)at tem-perature as low as400°C,as shown in Fig.1b.For sample sintered at500°C for4h,crystallinity improved further,but still no extra peaks other than those for BFO could be observed,as shown in Fig.1(S1).At high sintering temperature,nonper-ovskite phases such as Fe2O3and Bi2Fe4O9were observed,as shown in Fig.1(S2–S4).These extra phases are possibly due to volatization of Bi3+ionsThe Effect of Heat Treatment on Structural and Multiferroic Properties of Phase-Pure BiFeO32217and creation of oxygen vacancies and conversion of Fe 3+to Fe 2+ions.16Phase purity decreased from 100%to 86%from sample S1to S3and then increased slightly to 92%for sample S4,which is due to volatization of Bi 2Fe 4O 9impurity phase near 800°C.The lattice parameters and unit cell volume for all sintered samples were calculated for the hexagonal crystal structure using Chekcell soft-ware.An increasing trend in unit cell parameters aand volume V was observed from 5.57A˚to 5.61A ˚and 374A˚3to 378A ˚3for S1to S4,respectively,whereas the parameter c increased from 13.87A˚to 13.98A˚for S1to S3and then decreased to 13.87A ˚for S4.The ﬂuctuation in peak intensities as shown in the magniﬁed view of the XRD patterns for 2h (31°to 32.5°)in Fig.2conﬁrms the variation in the unit cell parameters.Shift of (104)and (110)peaks towards lower 2h angle indicates an increase in d spacing and strain in the unit cell.Volumetric strain calculated from unit cell parameters increased from 1.1910À3to 9.2910À3,which ispossibly due to the larger ionic radii of Fe 2+(0.77A˚)ions than Fe 3+(0.63A˚)ions.In Fig.2(S4),reduction in the intensity of peak (110)is a sign of structure modiﬁcation from rhombohedral to orthorhombic unit cell.17The crystallite size for all samples was calculated using the Scherrer formula:t ¼0:9k ðÞ=b cos #B ðÞðÞ;(1)where t is the crystallite size,k is the wavelength ofincident x-rays (k =1.5406A˚),#B is the diffraction angle,and b is the full-width at half-maximum (FWHM).Crystallite size was calculated for the major peak and showed an increasing trend (47nm to 67nm)for samples S1to S3and then decreased to 56nm for S4.Slight decrease in crystallite size is possibly due to modiﬁcation in crystallite shape and structure.Porosity of all the samples was deter-mined using the formula 18P ¼1Àq m =q t ðÞ;(2)where q m and q t are the measured and theoretical density,respectively.The densities (q m )were mea-sured using Archimedes’method,and the theoreti-cal density (q t =8.41g/cm 3)was calculated using the following formula:q t ¼nM ðÞ=VN A ðÞ;(3)where n is the number of formula units in the BFO hexagonal unit,M is the molecular mass of one formula unit,V is the volume of the unit cell,and N A is Avogadro’s number.The porosity of the sam-ples decreased from 0.56to 0.10after sintering.Microstructure AnalysisThe grain size,its distribution,and the surface morphology of the sintered pellets were observed by SEM.Figure 3shows SEM micrographs of pellets (S1–S4)sintered at different temperatures.It is observed that the sintering temperature signiﬁ-cantly affects the grain morphology and its growth.Grain size increased with sintering temperature,resulting in better ferroelectric connectivity;asaFig.1.XRD patterns of BiFeO 3for (a)as-synthesized powder,(b)powder calcined at 400°C in air for 3h,and for pellets then sintered at (S1)500°C,(S2)600°C,(S3)700°C,and (S4)800°C in air for 4h.Peaks are labeled with JCPDS cards no.01-072-7678(BiFeO 3),00-001-1053(Fe 2O 3),and 00-020-0836(Bi 2Fe 4O 9).Fig.2.Magniﬁed view of XRD pattern for 2h angle (31ºto 32.5º),showing the variation in (104)and (110)peaks of BiFeO 3for pellets sintered at (S1)500°C,(S2)600°C,(S3)700°C,and (S4)800°C.Shami,Awan,and Anis-ur-Rehman2218result,the density of the pellets improved and porosity decreased.The micrographs of sintered pellets show polygon morphology having rounded edges.This is an indication of melt-like behavior,just as in liquid-phase sintering.Thermal AnalysisThe thermal stability,phase transition,and melting temperatures of sintered BFO samples were determined by high-temperature DSC study.Figure 4shows DSC thermograms of BFO samples sintered at 500°C to 800°C.Two adjacent sharpendothermic peaks around 790°C and 820°C appeared for all samples,showing decreasing enthalpy with increasing sintering temperature.These can be interpreted as rhombohedral (a )to orthorhombic (b )and orthorhombic (b )to cubic (d )phase transitions.19These peaks are relatively clo-ser in our samples than previously reported for crushed dendritic single crystal,19and this variation is likely due to nanocrystallites in our BFO samples.The second transition corresponds to the ferroelec-tric to paraelectric Curie phase transformation temperature (T C ).20Finally,there is a sharp endo-thermic peak at 965°C,which corresponds to the heat loss at the melting point of BFO ceramics,in good agreement with literature.21Electrical DC Resistivity Measurements Electrical DC resistivity (q )measurements were carried out using the two-probe method 22,23from room temperature to 300°C as shown in Fig.5.Pressure contacts equal to the pellet size were used after polishing the surfaces.Resistivity at room temperature was in the range of G X -cm,declining to the M X -cm range with increasing sintering temper-ature.This depression is possibly due to reduction in porosity,well-connected grain boundaries,and for-mation of conduction ions at high sintering tem-peratures.All samples manifestedsemiconductingFig.3.SEM micrographs of BiFeO 3for pellets sintered at (S1)500°C,(S2)600°C,(S3)700°C,and (S4)800°C.Fig.4.DSC thermogram of BiFeO 3for pellets S1,S2,S3,and S4sintered at 500°C,600°C,700°C,and 800°C,respectively.The Effect of Heat Treatment on Structural and Multiferroic Properties of Phase-Pure BiFeO 32219behavior within the measuring temperature range.This semiconducting behavior of the samples may be due to the transport mechanism of polaron hopping,electrons thermally activated to the conduction band,and thermally assisted tunneling of charge carriers through grain boundaries.23Electrical AC PropertiesThe dielectric loss tangent [tan(d )],dielectric constant ( e ),and AC conductivity (r AC )of the sin-tered (BFO)sample were measured in the frequency range from 20Hz to 3MHz as shown in Fig.6.Measurements were conducted at room tempera-ture by parallel-plate technique.Two dielectric relaxation peaks were observed in the dielectric loss tangent at the extremes of the frequency measuring range,as shown in Fig.6a.At lower frequencies the contributions to the dielectric constant are due to space charges,electric dipoles,ions,andelectrons.Fig.5.Plot of electrical DC resistivity for pellets S1,S2,S3,and S4sintered at 500°C,600°C,700°C,and 800°C,respectively.Fig.6.Graphs of electrical AC measurements as a function of frequency from 20Hz to 3MHz on a log scale for sintered pellets S1,S2,S3,and S4showing (a)dielectric loss tangent [tan(d )],(b)dielectric constant e ,and (c)AC conductivity r AC .Shami,Awan,and Anis-ur-Rehman2220As the frequency is increased from mHz to a few Hz,space charges do not have sufﬁcient time to build up interfacially and undergo the ﬁrst relaxation.At high frequencies,the second relaxation is due to dipole moments,which do not have sufﬁcient time to align along the applied varying electric ﬁeld,and their contribution dies out.24The dielectric constant ( e )was calculated by the following relation:e ¼Cd ðÞ=e 0A ðÞ;(4)where d is the pellet thickness,C is the capacitance,e 0is the permittivity of free space,and A is the cross-sectional area of the pellet.Sample S4in Fig.6b exhibited the maximum value of dielectric constant,which is seemingly due to high density,well-connected grain boundaries,and excess of free ions.All samples displayed a declining progression with increasing frequency.Step-like reduction in dielectric constant indicates lapse of interfacial polarization and then dipole polarization.This dielectric behavior can be explained based on the Maxwell–Wagner model.25The AC conductivity was also evaluated from the dielectric constant ( e )and dielectric loss tangent [tan(d )]in the same frequency range mentioned above using the following relation:18r AC ¼xe 0 e tan d ðÞ;(5)where r AC is the AC conductivity and x is the angular frequency.It can be observed from Fig.6c that r AC increases with increasing frequency,which is in accordance with the Maxwell–Wagner model.25The increase in the electrical conductivity with increasing sintering temperature can be explained by the Koops model.26According to this model,polycrystalline materials have well-conducting grains of resistivity q g and thickness t g which are separated by poorly conducting grain boundaries of resistivity q gb and thickness t gb .The relation of resistivity for such materials will beq s ¼q g þt gb =t g ÀÁq gb :(5)Increasing the sintering temperature will shrinkthe grain boundary thickness and enlarge the grain thickness.Therefore,a decrease of (t gb /t g )will result in an increase of the electrical AC conductivity of the sample.Magnetic CharacterizationThe magnetization as a function of applied mag-netic ﬁeld (±10kOe)at room temperature for all samples is shown in Fig.7,whereas the inset shows a magniﬁed view of selected area at low applied ﬁeld (±150Oe).The magnetization versus applied mag-netic ﬁeld (M –H )loops exhibit almost linear dependence on magnetization ﬁeld.Although BFO exhibits antiferromagnetic nature in bulk samples,weak ferromagnetic features can be seen in all the samples here.27With the decrease in particle size,the surface-to-volume ratio becomes large and long-range antiferromagnetic order is frequently interrupted at the surfaces.The contribution of uncompensated spins at the surface to the total magnetic moment of the particle increases.An intrinsic spiral spin structure is partially suppressed,which leads to weak ferromagnetism of the nano-particles.28Maximum remnant magnetization (M r =1.9910À3emu/g)and maximum coercivity (H c =120.3Oe)were observed for sample S3.This increase in magnetization may be due to two factors related to increasing sintering temperature.One is the creation of optimum ratio of Fe 3+to Fe 2+ions after incorpora-tion of impurity phases,and the second is strain in the structure which has diverted the canting angle of magnetic moments in cycloidal order.FerroelectricityThe ferroelectric behavior of the samples was investigated at room temperature by polarization versus applied electric ﬁeld (PE)loops as shown in Fig.8a.PE loops were not saturated for all the sam-ples,which may be due to leakage currents,as BFO bulk sample showed conductive behavior at high voltage.29Variation in shape of the loops,remnant polarization and coercivity were observed with increasing sintering temperature.Lossy capacitor behavior in PE loops increases with increasing sin-tering temperature.Irrespective of the relaxation peaks,similar trend can also be observed in Fig.6a for tan(d ),where the dielectric loss tangent has an increasing trend with increasing sintering tempera-ture.These tendencies are possibly due to rise in leakage current at elevated sintering temperatures.Remnant polarization (2P r )and coercive ﬁeld (2E c )were measured from the PE loops and are plotted in Fig.8b.Remnant polarization showed an increasing trend with increasing sintering temper-ature,and sample S3exhibited maximum valueofFig.7.M –H loops for BiFeO 3pellets S1,S2,S3,and S4sintered at500°C,600°C,700°C,and 800°C,respectively.Inset shows partially enlarged view from the selected area.The Effect of Heat Treatment on Structural and Multiferroic Properties of Phase-Pure BiFeO 32221Fig.8.PE measurement results of sintered pellets S1,S2,S3,and S4:(a)for ±100kV/cm,and (b)variation in remnant polarization (2P r )and coercive ﬁeld (2E c )with sinteringtemperature.Fig.9.Plots of capacitance as a function of temperature from 30°C to 400°C at constant frequencies of 3MHz,2MHz,1MHz,and 0.5MHz for BiFeO 3pellets sintered at (S1)500°C,(S2)600°C,(S3)700°C,and (S4)800°C.Shami,Awan,and Anis-ur-Rehman22222P r =2.78l C/cm 2.The coercive ﬁeld exhibits a sinusoidal trend,which is possibly due to variation in volumetric strain in the unit cell,as shown in Fig.2.Maximum coercive ﬁeld 2E c =116kV/cm 2was also observed for sample S3.The enhanced remnant polarization and coercive ﬁeld in S3rela-tive to the other samples is due to the optimum tilting of oxygen octahedra with respect to the bis-muth ions in the BFO perovskite structure.Magnetoelectric EffectThe magnetoelectric (ME)effect was observed in two different ways:ﬁrstly by measuring the capac-itance as a function of temperature from 30°C to 400°C at constant frequencies of 0.5MHz,1MHz,2MHz,and 3MHz,and secondly by measuring the capacitance of the samples in the frequency range from 20Hz to 3MHz with and without magneticﬁeld.Figure 9shows the capacitance as a function of temperature for all the samples.Maxima incapacitance can be observed around the Ne´el tem-perature (T N )for all the samples.14With increasing temperature,electric dipoles acquire more favorable conditions to orient in the direction of ferromagneticmoments within the sample.Above the Ne´el tem-perature,the ME effect starts to disappear,and the electric dipoles become disordered due to the para-magnetic behavior of the samples,and ultimately the capacitance of the samples exhibits a decreasing trend.T N is different for different samples but comparable to literature values,30shifting towards lower temperature with increasing sintering tem-perature.The capacitance peak for samples S1to S3is at the same temperature for the four applied frequencies.However,the case is different for sample S4,where the capacitance peaks shift with the applied frequency.This implies thattheFig.10.Graphs of capacitance as a function of frequency from 20Hz to 3MHz on a log scale with and without magnetic ﬁeld for BiFeO 3sintered pellets (a)S1,(b)S2,(c)S3,and (d)S4.Insets of (c)and (d)show magniﬁed views of selected area from their corresponding graphs.The Effect of Heat Treatment on Structural and Multiferroic Properties of Phase-Pure BiFeO 32223antiferromagnetic order in S4is very weak and activation energy from the applied electricﬁeld can alter the Ne´el temperature.Figure10showed the variation in capacitance with and without application of an external applied magneticﬁeld.In the samples without any external ﬁeld,the electric dipole moments try to align in the direction of magnetic moments,which are in ferro-magnetic order.As the external magneticﬁeld is applied,the magnetization of the sample increases, which means that the magnetic moment has tried to align in the direction of the appliedﬁeld.Ulti-mately,the electric dipoles also reorient in the direction of magnetic moments,which increases the capacitance of the material.This effect can be observed for all the samples and seems to decrease with increasing sintering temperature for sam-ples S3and S4,although this apparent effect is due to larger capacitance values in these samples.Insets of Fig.10present magniﬁed views for S3and S4, showing the variation in capacitance due to the applied magneticﬁeld.CONCLUSIONSHighly resistive,phase-pure BiFeO3has been prepared using a diverse precursor bismuth oxide at temperature as low as400°C.Sintering tempera-ture above500°C signiﬁcantly inﬂuenced the structural and multiferroic properties.XRD analy-sis conﬁrmed the perovskite structure of the sam-ples and showed that the crystallite size increased with increasing sintering temperature.SEM micrographs showed grain growth during sintering and spherical morphology of grains,showing melt-like behavior.DSC analysis showed that the Curie temperature of pure BiFeO3is820°C±1°C.The sintering temperature signiﬁcantly inﬂuenced the dielectric behavior as a function of frequency, showing two relaxation peaks for dielectric disper-sion.For the resistivity as a function of tempera-ture,all the samples showed a semiconducting trend following Arrhenius behavior.Magnetization showed linearﬁeld dependence behavior at room temperature.PE measurements provided evidence of the ferroelectric nature of all the samples,while maximum remnant polarization and coercivity were observed for sample S3.Capacitance as a function of temperature showed maxima near the Ne´el tem-perature,which shifted to higher temperature for higher applied frequency.Change in capacitance under applied magneticﬁeld showed the magneto-electric response.REFERENCES1. C.H.Wang,Z.F.Liu,L.Yu,Z.M.Tian,and S.L.Yuan,Mater.Sci.Eng.B176,1243(2011).2.V.S.Puli, A.Kumar,N.Panwar,I.C.Panwar,and R.S.Katiyar,J.Alloys Compd.509,8223(2011).3.Z.Wen,L.You,X.Shen,X.Li,D.Wu,J.Wang,and A.Li,Mater.Sci.Eng.B176,990(2011).4.G.Biasotto,A.Z.Simo˜es,C.R.Foschini,M.A.Zaghete,J.A.Varela,and E.Longo,Mater.Res.Bull.46,2543(2011). 5.K.Liu,H.Fan,P.Ren,and C.Yang,J.Alloys Compd.509,1901(2011).6. D.Varshney,A.Kumar,and K.Verma,J.Alloys Compd.509,8421(2011).7.X.Zhang,Y.Sui,X.Wang,J.Mao,R.Zhu,Y.Wang,Z.Wang,Y.Liu,and W.Liu,J.Alloys Compd.509,5908 (2011).8. 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热处理对BT36高温钛合金组织及性能的影响

第23卷第2期2003年6月航空材料学报JOURNAL OF AERONAUT ICAL MATERIALS Vol.23No.2June.2003热处理对BT36高温钛合金组织及性能的影响郝孟一,蔡建明,杜娟(北京航空材料研究院,北京100095)摘要:B T36是俄罗斯研制的一种600e 高温钛合金,该合金的特点之一是含5%左右的W 。

本文采用B T36轧棒(<20mm)研究了5种不同的热处理制度,进行了金相组织分析和性能测试(包括拉伸、蠕变、热稳定性等)。

结果表明,获得片层组织是提高合金600e 持久及蠕变性能的关键,而且需要依靠适当的热处理工艺控制片层的厚度及形貌。

B T 36棒材通过5重热处理(995e P 0.5h,AC+970e P 2h,FC y 700e P 2h,AC +500e P 15h,AC+650e P 6h,AC)可获得较好的综合性能。

关键词:B T 36高温钛合金;热处理;显微组织;力学性能中图分类号:TG146.2+3 文献标识码:A 文章编号:1005-5053(2003)02-0014-04600e 高温钛合金是制造高性能航空发动机的重要结构材料之一,主要用于压气机段的静子叶片、转子叶片、轮盘件等。

目前,典型的600e合金有英国的I MI834[1](T-i 5.8A-l 4Sn -3.5Zr -0.7Nb -0.5Mo -0.35S-i 0.06C),美国的T-i 1100[2](T-i 6A-l 2.75Sn -4Zr -0.4Mo -0.45Si),俄罗斯的B T36[3](T-i 6.2A-l 2Sn -3.6Zr -0.7Mo -5.0W -0.15S-i 0.1Y)和中国的T-i 60(T-i 5.6A-l 4.8Sn -2Zr -1Mo -0.35S-i 1Nd)。

以上这些合金均为T-i A-l Sn -Zr -Mo -Si 合金系。

时效温度对Super304H钢析出相的影响(1)

第30卷第6期2009年 12月材料热处理学报TRANS AC TIONS OF MATERIALS AND HEAT TREATMENTVo l .30 No.6December 2009时效温度对Super304H 钢析出相的影响李新梅1, 邹勇1, 张忠文2, 邹增大1(1.山东大学材料液固结构演变与加工教育部重点实验室,山东济南 250061;2.山东电力研究院,山东济南 250002)摘要:利用扫描电子显微镜、电子探针、X 射线衍射和透射电镜研究了新型奥氏体耐热钢Super304H 在高温时效条件下析出相的变化。

结果表明,Super304H 钢经700~1250e 时效后,组织中出现4种析出相:N b(C,N)、富Cu 相、M 23C 6和NbCrN 。

随时效温度的不同,析出相发生析出或溶解的变化,同时它们的形态、分布和数量随温度变化呈现出不同的变化规律,其中M 23C 6在700~900e 主要沿晶界析出,这将会降低钢的高温蠕变强度及抗晶间腐蚀性能。

关键词:Super304H 钢; 时效; 析出相中图分类号:TG14217 文献标识码:A 文章编号:100926264(2009)0620051206Effect of aging temperature on precipitation of Super 304H steelLI Xin 2mei 1, Z O U Yong 1, Z HA NG Z hong 2wen 2, Z O U Zeng 2da1(1.Key Laboratory of Liquid 2Solid S truc tural Evolution and Processing of M aterials(M inistry of Education),Shandong University,Jinan 250061,China; 2.Shandong Electric Power Research Institute,Jinan 250002,China)Abstr act :Microstructure o f no vel austenitic heat 2resistance steel Super304H after aging at different temperatures w as investigated by means of scanning electron micro scopy (SE M),electron probe microanalysis (EPM A),X 2ray diffraction (X RD)and transmissio n electro n microsco py (TEM).The experimental results indicate that the precipitated phases o f Nb(C,N)、Cu 2rich phase,M 23C 6and N bCrN are detected depending on different aging temperatures,when the super304H steel is aged at 700~1250e .Wi th the v ariation o f aging temperature,the different precipitated phases were observ ed i n the steel,and the mo rpholog y and amo unt of the precipitates changed.M 23C 6carbides mainly precipitated along grain bo undary for the steel aged at 700~900e ,leading to deteriorate high temperature creep strength and interg ranular corrosion resistant of Super304H steel.Key wor ds :Super304H steel;aging treatment;precipitation收稿日期: 2008210217; 修订日期: 2009204214基金项目: 山东省电力集团公司重点科技项目(2007A 247);山东省自然科学基金(Z2006F07)作者简介: 李新梅(1980)),女,山东大学博士,研究方向:奥氏体耐热钢,发表论文7篇,电话:0531288396145,E 2m ail:li xinm ei128@163.c om 。

Influence of heat treatment conditions on the structure and magnetic properties of barium ferrite Ba

Materials Chemistry and Physics98(2006)90–94Inﬂuence of heat treatment conditions on the structure and magnetic properties of barium ferrite BaFe12O19hollow microspheres of low densityPing Ren a,b,JianGuo Guan b,∗,XuDong Cheng ba Department of Energy Sources and Environment Engineering,Shanghai University of Electric Power,Shanghai200093,PR Chinab State Key Laboratory of Advanced Technology for Materials Synthesis,Wuhan University of Technology,Wuhan430070,PR ChinaReceived22May2005;received in revised form28August2005;accepted28August2005AbstractBarium ferrite BaFe12O19hollow microspheres of low density(about2.50g cm−3)were synthesized without BaFe2O4or other intermediate phase by spray pyrolysis technique,combined with co-precipitation precursor and followed further heat treatment.The relationships between the microstructure,magnetic properties and the sintering temperature,time were investigated.Hollow microspheres,constituted by nanometer particles,showed a broad particle distribution of2–15␮m.When the sample was sintered at900◦C for3h or at1100◦C for2h,pure BaFe12O19 ferrite was formed.With the increase of the sintering temperature or time,the grain size and preferential growth orientation were inﬂuenced,the saturation magnetization(M s)had a small increase,and the coercive force(H c)decreased obviously.©2005Elsevier B.V.All rights reserved.Keywords:Barium ferrite;Hollow microspheres;Spray pyrolysis;Microstructure;Magnetic properties1.IntroductionBarium ferrite(BaFe12O19)has been well-known for perma-nent magnets since its development by Philips researchers in the beginning of the1950s because of their high-coercivity,high-saturation magnetization,high resistance,low cost,associated with their excellent chemical stability and resistance to corrosion [1–3].In recent years,hexagonal ferrite has also caused wide interest in high-density recording media and radar-absorbing coatings[4,5].Its frequency of resonance absorption in the range of GHz is more than one thousand times as high as that of spinel ferrite.Therefore,it is promising for a kind of absorption mate-rial in theﬁeld of high frequency.However,they are quite heavy, which restricts their applications requiring lightweight mass. Moreover,they have difﬁculties in increasing the permeability in GHz region because of Snoek limit[6,7].As one of the ways to overcome these problems,hollow microspheres of low den-sity are suggested.At the same time hollow microspheres also exhibit different catalytic,optical,electric and magnetic prop-erties and have important applications in materials,chemistry, catalysis and biologyﬁeld[8].It is expected that microwave ∗Corresponding author.Tel.:+862787218832;fax:+862787879468.E-mail address:guanjg@(J.G.Guan).absorbing ability can be improved by adjusting the diameter and wall thickness of hollow microspheres or composition and grain size of ferrite nanometer particles.A number of fabrication methods have been used to produce hollow microspheres which are comprised of polymer,metal and ceramic materials,including nozzle reactor approaches (spray drying or pyrolysis)[9,10],emulsion/phase separation techniques(often combined with sol–gel processing)[11], emulsion/interfacial polymerization strategies[12],deposition method on sacriﬁcial cores[13]and self-assembly process [14,15].Among these routes,spray pyrolysis technique was considered simplest and suitable for the synthesis of inorganic ferrite hollow microspheres,and large quantity production is likely by using the route.In this paper,we synthesized ferrite precursor with co-precipitation method and obtained BaFe12O19 hollow microspheres by spray pyrolysis technique followed fur-ther heat treatment,and reported the inﬂuence of heat treatment conditions on theﬁnal structure and magnetic properties of the samples.2.ExperimentalBarium ferrite hollow microspheres were produced by three steps,the inves-tigated variables were:calcination temperature(T)and calcination time(t). Firstly,an aqueous solution of Fe(NO3)3·6H2O and Ba(NO3)3·2H2O(Fe/Ba0254-0584/$–see front matter©2005Elsevier B.V.All rights reserved. doi:10.1016/j.matchemphys.2005.08.070P.Ren et al./Materials Chemistry and Physics98(2006)90–9491atomic ratio equal to12)was added to a stirring aqueous solution of NaOH and Na2CO3(OH−/CO32−ion ratio equal to5).After they reacted for2h,the resulting precipitates standed for24h,were then washed with deionized water and dried at110◦C.The dry precipitate was grinded in order to obtainﬁne and uniform barium ferrite precursor powder.Secondly,ﬁne barium ferrite precursor powder was atomized into small particles by negative pressure of the FS-4Flame Spray Dryer.At the same time they were heated quickly at the high temperature ﬂame,which was produced by acetylene and oxygen as combustion gas.They passed the different holes of the Flame Spray Dryer respectively,mixed and burnt in the nozzle.Therefore,spherical droplets,caused by surface tension, decomposed to obtain hollow microspheres.Finally,the hollow microspheres were colleted and further calcined to form BaFe12O19phase or promote crystal-lization under air for3h at several different temperatures(600,800,1000,1100, 1200◦C)and at1100◦C for different time(2,3,4h).X-ray diffraction with Cu K␣radiation(XRD)was used for phase analysis. The surface morphology and size of the grains were studied by scanning electron microscope(SEM).The thermal curves were presented by thermo gravimetric analysis(TG)and differential scanning calorimeter(DSC).The magnetic mea-surements are performed using a vibrating sample magnetometer(VSM)to an appliedﬁeld of16000Oe,hysteresis loops are measured at room temperature.3.Results and discussion3.1.Structural analysis of pyrolysis products without heat treatmentFig.1shows the micrograph,phase constitution and the thermal curves for the pyrolyzed sample by SEM,XRD,TG and DSC.Spherical particles with a broad particle distribu-tion of2–15␮m were observed in Fig.1(a).This was prob-ably related to uneven solid ferrite precursor powder from co-precipitation method or the uniform atomization of the equip-ment.It demonstrated microspheres could be formed by spray pyrolysis method.Although hollow structure could not directly be seen in Fig.1(a),the average density of the sample(about 2.50g cm−3),which represented a50%decrease of hollow microspheres compared to solid powder of(about5.31g cm−3), could explain it to some extent.A small number of slight diffrac-tion peaks␣-Fe2O3and BaCO3were observed in Fig.1(b), which indicated the sample was constituted by a mixture of␣-Fe2O3and BaCO3phase of low crystallinity and also implied time of pyrolysis was not long enough to completely decom-pose iron hydroxide and barium carbonate or to form BaFe12O19 phase.Therefore,its necessary for further heat-treatment to form ferrite phase.Fig.1(c)was evident,an endothermic reaction accompanying with5.2%weight loss was observed at the tem-perature between78and138◦C,which showed the precursor did not decompose completely during the pyrolysis process. Between838and938◦C,the sample undergone exothermic reaction without weight loss.At the temperature range by com-paring XRD patterns of the samples after being heated at600 and1200◦C,it revealed that this change could be ascribed to the reaction of oxides to form barium ferrite.3.2.Inﬂuence of the heating temperature on the structureand magnetic propertiesThe phase constitution of the samples heat-treated at differ-ent temperatures for3h with a heating rate of5K min−1was obtained by XRD and is shown in Fig.2.At600◦C,a small number of␣-Fe2O3diffraction peaks and some slight peaks of BaFe12O19phase were observed.Single phase barium fer-rite had started to form by800◦C(except for the only slight peak at2.69˚A which corresponded to␣-Fe2O3).At900◦C,the sample presented a series of peaks all assigned to BaFe12O19 Fig.1.(a)SEM image,(b)XRD pattern and(c)TG and DSC curves of the sample after spray pyrolysis without heat treatment.92P.Ren et al./Materials Chemistry and Physics98(2006)90–94Fig.2.XRD patterns of the samples calcined at different temperature for3h: (a)600◦C,(b)800◦C,(c)900◦C,(d)1100◦C,(e)1200◦C.phase.When the temperature increased further,the values of diffraction peak intensities improved obviously,which could be associated with the amount and grain coarsening of BaFe12O19 phase.There was a slight change between theﬁrst and second strongest peaks at1200◦C due to preferential grain growth.In this work,pure barium ferrite BaFe12O19was formed directly from iron oxides and barium oxides as the onlycomponents Fig.4.Static magnetic properties of the samples calcined at different tempera-tures for3h.detected.Other intermediate phase such as␣-BaFe2O4,␥-Fe2O3 were never observed during the heat treatment.This is consistent with the results obtained in the previous work[16].As expected from XRD data,change in grain morphology with temperature was observed,which is presented in Fig.3. Hollow microspheres had kept intact during the calcination pro-cess.The surface of the sample without heat treatment was relatively smooth.The surface became rougher with the increase of sintering temperature.By1100◦C,the spheres were consti-tuted by crystals for hexagonal like platelet of about several hundred nanometer.These clearly indicated heat treatment was beneﬁcial for ferrite formation and crystallization.Fig.4shows the effect of the sintering temperature on the magnetic properties between600and1200◦C for3h.The sat-uration magnetization(M s)increased obviously from600to 1000◦C,increased slightly between1000and1200◦C,and reached a maximum of55emu g−1at1200◦C.The value is lower than the theoretical limit(72emu g−1)[17],which might be associated with the density and porosity of hollow micro-spheres.A drastic increase of M s between800and1000◦Cwas Fig.3.SEM images of the samples calcined at different temperatures for3h:(a)without heat treatment,(b)800◦C,(c)1000◦C,(d)1100◦C.P .Ren et al./Materials Chemistry and Physics 98(2006)90–9493related to the increasing amount of BaFe 12O 19phase and a small increase from 1000to 1200◦C for grain coarsening according to XRD and SEM results.The hematite formation in the samples reduced the volumetric fraction of barium ferrite and conse-quently reduced M s before 800◦C.This result is in accordance with that previously described in literature [18].The rema-nent magnetization (M r )reached a maximum of 28emu g −1at 1000◦C,decreasing slightly for higher temperature.The intrin-sic coercive force (H c )basically showed a contrary behavior as M s ,increased obviously before 800◦C due to the increas-ing amount of BaFe 12O 19phase,decreased slightly from 3.4to 3.2kOe between 800and 1000◦C which is close to the coer-civity value of typical hard magnetic materials,and decreased drastically from 3.2to 1.3kOe between 1000and 1200◦C owing to grain coarsening.This indicated the coercive force exhibited strong dependency on the sintering temperature.Similar results were reported by Joonghoe et al.[19],who studied effects of the grain boundary on the coercivity of barium ferrite BaFe 12O 19.The results were also supported by Janasi et al.[18]who con-ducted an investigation on the effect of heat treatment conditions and routes on the magnetic properties of co-precipitated ferrite powders.3.3.Inﬂuence of the heating time on the structure and magnetic propertiesFig.5shows XRD patterns of the samples calcined at 1100◦C for different heat treatment times.In all cases,the samples were virtually single phase BaFe 12O 19ferrite.With the extending of sintering time,the relative intensities of diffraction peaks increased.At the same time the average grain size,measured from the full-width at half-maximum (FWHM)of (107)peak using Scherrer formula,was found to increase from about 450nm for the heated powder for 2h,520nm for 3h,to 550nm for 4h.According to the XRD intensity formulaI =I 0λ3e 432πR 3m 2c 4 Vq V 20|F h k l |2 1+cos 22θsin 2θcos θFig.5.XRD patterns of the samples calcined at 1100◦C for different times:(a)2h,(b)3h,(c)4h.Fig.6.Hysteresis curve of the samples calcined at 1100◦C for different time.where V is the grain volume of (h k l )plain-oriented,q thegrain number of (h k l )plain-oriented,V 0the volume of unit cell,the relative intensities of diffraction peaks were consis-tent with the calculated intensities for the powder heat treated at 1100◦C for 3h.However,when the sintering time was 4h,(107)diffraction peak were dominant whereas relative intensity of (114)diffraction peak was diminished because the grain size of (107)plain-oriented increased quicker than that of (114)plain-oriented based on the above formula and Scherrer for-mula.These implicated that preferential grain growth probably occurred when the samples were sintered at long time.Fig.6shows the magnetization hysteresis loops at room tem-perature of the samples sintered at 1100◦C between 2and 4h.All the samples had almost identical saturation magnetization and remanent magnetization values,indicating that M s and M r are independent on the sintering time,because the samples were all pure BaFe 12O 19phase and their densities are basically same.On the other hand,the coercivity was changed with the sinter-ing time,which decreased initially from 3.0kOe for 2h and then recovered to 2.5kOe for 4h and reached the minimum of 1.9kOe for 3h because of grain size and preferential grain growth according to the XRD results in Fig.5.This clearly illus-trated the well-known fact that the coercivity strongly depended on the calcining time.4.ConclusionsWe had prepared barium ferrite BaFe 12O 19hollow micro-spheres mainly distributed between 2and 15␮m by spray pyrolysis technique,combined with co-precipitation precursor and followed further calcinations,and studied the effect of heat treatment conditions.We found that hollow microspheres kept unbroken and the average density nearly kept unchange-able (about 2.50g cm −3)during the heat treatment.When the sample was sintered at 900◦C for 3h or at 1100◦C for 2h,the single phase BaFe 12O 19ferrite was formed,directly by iron oxide and barium oxide without BaFe 2O 4or ␥-Fe 2O 3intermediate phase.Increasing the temperature was beneﬁcial in increasing the crystalline,improving the saturation mag-netization and eliminating the coercive force.The saturation94P.Ren et al./Materials Chemistry and Physics98(2006)90–94magnetization(M s)increased obviously before1000◦C,and increased slightly between1000and1200◦C and reached the maximum of55emu g−1at1200◦C.Increasing the heat treat-ment time inﬂuenced grain size,preferential grain growth ori-entation and magnetic properties.H c reached a minimum of 1.9kOe for3h and a maximum of3.0kOe for2h,whereas M s was little affected.In a conclusive manner we can say that H c was strongly dependent on heat treatment conditions owing to grain size and preferential growth and M s was closely related to the materials composition.Preparation of narrow distribution of smaller hollow microspheres and study of their structure and properties will be our following work.AcknowledgementThis work is supported by863HI-TECH project of China (no.2002AA305302).References[1]J.Smit,H.P.J.Wijn,Ferrites,Philips Technical Library,Eindhoven,1959.[2]T.Fujiwara,IEEE Trans.Magn.21(1985)1480.[3]P.Campbell,Permanent Magnet Materials and their Application,Cam-bridge University Press,Cambridge,1994.[4]N.Ichnose,Super Particle Technology,Springer-Verlag,1992.[5]K.J.Viney,R.M.Jha,Radar Absorbing Materials,Kluwer AcademicPublishers,1996.[6]K.Hatakeyama,T.Inui,IEEE Trans.Magn.20(1984)1261.[7]M.Matsumoto,Y.Miyata,IEEE Trans.Magn.33(1994)4459.[8]D.L.Wilcox,T.Bernat,D.Kellerman,J.K.Cochran,Materials ResearchSociety Proceedings,vol.375,Pittsburgh,1995.[9]M.Iida,T.Sasaki,M.Watanabe,Chem.Mater.10(1998)3780.[10]K´a roly Zolt´a n,Sz´e pv¨o lgyi J´a nos,Powder Technol.132(2/3)(2003)211.[11]T.Z.Ren,Z.Y.Yuan, B.L.Su,Chem.Phys.Lett.374(1/2)(2003)170.[12]J.Hotz,W.Meier,Langmuir14(1998)1031.[13]J.L.Yin,X.F.Qian,J.Yin,M.W.Shi,G.T.Zhou,Mater.Lett.57(24/25)(2003)3859.[14]R.A.Caruso,A.Susha,F.Caruso,Chem.Mater.13(2001)400.[15]F.Caruso,R.A.Caruso,H.Mohwald,Science282(6)(1998)1111.[16]R.C.Pullar,A.K.Bhattacharya,Mater.Lett.57(2002)537–542.[17]B.T.Shirk,W.R.Buessem,J.Appl.Phys.40(1969)1294.[18]S.R.Janasi,M.Emura,ndgraf,D.Rodrigues,J.Magn.Magn.Mater.238(2002)168–172.[19]Joonghoe Dho,E.K.Lee,J.Y.Park,N.H.Hur,J.Magn.Magn.Mater.285(2005)164–168.。

Effects of heat treatments on the microstructure and mechanical properties of a 6061 aluminium alloy

Materials Science and Engineering A 528 (2011) 2718–2724Contents lists available at ScienceDirectMaterials Science and EngineeringAj o u r n a l h o m e p a g e :w w w.e l s e v i e r.c o m /l o c a t e /m s eaEffects of heat treatments on the microstructure and mechanical properties of a 6061aluminium alloyD.Maisonnette a ,M.Suery b ,D.Nelias a ,∗,P.Chaudet a ,T.Epicier caUniversitéde Lyon,CNRS,INSA-Lyon,LaMCoS UMR5259,F-69621,FrancebUniversitéde Grenoble,SIMaP,UMR CNRS 5266,BP46,Domaine Universitaire,38402Saint Martin d’Hères Cedex,France cUniversitéde Lyon,CNRS,INSA-Lyon,Mateis UMR5510,F-69621,Francea r t i c l e i n f o Article history:Received 23August 2010Received in revised form 3December 2010Accepted 3December 2010Available online 9 December 2010Keywords:6061Aluminium alloyThermomechanical propertiesElectron beam welding stress–strain curves Yield stressHardening precipitatesa b s t r a c tThis paper describes the mechanical behavior of the 6061-T6aluminium alloy at room temperature for various previous thermal histories representative of an electron beam welding.A fast-heating device has been designed to control and apply thermal loadings on tensile specimens.Tensile tests show that the yield stress at ambient temperature decreases if the maximum temperature reached increases or if the heating rate decreases.This variation of the mechanical properties is the result of microstructural changes which have been observed by Transmission Electron Microscopy (TEM).© 2010 Elsevier B.V. All rights reserved.1.IntroductionThe study presented in this paper is concerned with the widely used 6061-T6aluminium alloy.It is an age hardenable alloy,the mechanical properties of which being mainly controlled by the hardening precipitates contained in the material.When the material is subjected to a solution heat treatment followed by a quenching and a tempering treatment,its mechanical properties reach their highest level and become very good compared to other aluminium alloys.The as-obtained microstructure of the material is called T6temper (tempering around 175◦C).Another interest-ing characteristic of the AA6061is its good weldability.Because of these favorable properties,the AA6061alloy is used in the trans-port and the public works domains (framework,pylon,handling equipment ...)and also for complex structures assembled by weld-ing [1–3].The present work is part of the early qualifying study of a pres-sure vessel to be used in an experimental nuclear reactor.The approximate size of the vessel is ﬁve meters height with a diameter of about one meter.Several ferrules in AA6061-T6should be assem-bled together by electron beam (EB)welding.The aim of the work presented in this paper is to evaluate the inﬂuence of the weld-ing process on the mechanical properties of the material at room∗Corresponding author.E-mail address:daniel.nelias@insa-lyon.fr (D.Nelias).temperature.The change of mechanical properties is due to met-allurgical phenomena such as dissolution,growth or coarsening of precipitates,which have been also observed.It is commonly assumed that the generic precipitation sequence in Al–Mg–Si alloys is [4,5]:SSSS →GP →␤→␤→␤-Mg 2Si(1)Here SSSS represents the super-saturated solid solution and GP stands for Guinier–Preston zones.The sequence (1)will be consid-ered in this work.However,some authors give more details about this sequence [5–12]particularly Ravi and Wolverton [5]who gave a detailed inventory of the compositions of the phases contained in an Al–Mg–Si alloy.The compositions generally accepted for the most common precipitates are listed in Table 1.According to the literature [6–9,13,14],the T6temper of the 6XXX alloys involves very thin precipitates.They are ␤ needle-shaped precipitates oriented along the three 100 directions of the matrix.Their size is nanometric and they are partially coherent.The study presented in this paper includes High Resolution Transmission Electron Microscopy (HRTEM)observations of the investigated 6061-T6alloy in order to characterize the precipita-tion state of the T6temper.These observations will allow deﬁning a precipitate distribution of reference for the initial alloy.From this initial state,thermal loadings are applied on specimens which are thereafter observed by TEM.The investigated thermal loadings will also be applied on tensile specimens in order to evaluate the variation of the resulting mechanical properties.0921-5093/$–see front matter © 2010 Elsevier B.V. All rights reserved.doi:10.1016/j.msea.2010.12.011D.Maisonnette et al./Materials Science and Engineering A528 (2011) 2718–27242719Table1Compositions of the precipitates contained in Al–Mg–Si alloys.Phase CompositionGP zone Mg1Si1␤ Mg5Si6␤ Mg9Si5␤Mg2SiFor experimental convenience,the study will be limited to the solid state of the alloy.This means that the maximum temperature to be used is below582◦C(solidus temperature for the AA6061)and the phenomena occurring in the melting pool of the weld will not be taken into account here.Furthermore,the mechanical characteriza-tions and microstructural observations will be carried out at room temperature after the thermal loading.This will allow the char-acterization of the material at various points of the Heat Affected Zone(HAZ)after welding(and not during the welding process). For that purpose,the required thermal loadings should reproduce the temperature evolution in the HAZ with high heating rates up to200K/s.An experimental device has been speciﬁcally developed to meet these requirements.Atﬁrst,the design of the device will be brieﬂy presented.Then,the results of the mechanical charac-terizations and microstructural observations will be presented and discussed.2.Experimental procedure2.1.Experimental heating deviceThe main purpose of the experimental heating device is to repro-duce on a tensile specimen the thermal history encountered by each point of the heat affected zone during welding of the vessel.The highest temperature to be studied is thus T=560◦C,very close to the solidus temperature of582◦C which should not be reached.To do so,an accurate control of the temperature has been set up.Fur-thermore,the device should be able to reproduce the heating rate observed in the HAZ of an electron beam welding(up to200K/s). This heating rate has been evaluated by measuring it during an instrumented welding experiment.The second aim of the device is to apply a mechanical loading on a specimen in order to mea-sure the mechanical properties of the material.The mechanical and thermal loadings have to be used simultaneously in order to perform tensile tests at high temperature for further study or to compensate for thermal expansion of the specimen during heating. Therefore,the experimental equipment includes a heating device and a mechanical testing machine.2.1.1.Design of the deviceA convenient method to heat aluminium alloys at very high rate is by Joule effect.Another way would be by induction heating but it is not efﬁcient enough to obtain the required heating rate on alu-minium alloys.For this reason,a resistive heating device has been designed and constructed.In order to measure the temperature of the specimen during heating,a thermocouple has been spot welded on the specimen surface.The strains are measured by means of an extensometer with ceramic tips.The Joule heating device is a power supply,made of an electrical transformer and a thyristor bridge,providing a continuous current whose intensity is controlled by a thermal controller.Water cooled cables and clamping systems are used to connect the specimen to the heating device.A graphite resistor is added in series in order to increase the potential difference across the generator allowing a good temperaturecontrol.Fig.1.Temperature distribution measured by thermocouples along the tensile specimen.2.1.2.Specimen designA specimen heated by using Joule effect reacts as an electrical resistor.Its electrical resistance depends on the material electri-cal resistivity and the specimen shape which has to be optimized in order to reach the desired heating rate(up to500K/s).More-over,the temperature must be uniform over the measurement area (between the extensometer tips)and the specimen volume should be large enough for the microstructure to be representative of the alloy in real structures.A FEM simulation was performed to optimize the size and shape of the specimen.The used software,called Sysweld®was devel-oped by ESI Group.The simulation is carried out by using an electro kinetic model[15].The density d and the thermal conductivity K of the alloy were considered to vary with temperature.A paramet-ric study shows that a diameter of6mm is required to obtain a heating rate up to500K/s.A specimen length of100mm is also required to have a low thermal gradient.Fig.1shows the tempera-ture distribution in the specimen.The gradient has been measured with10thermocouples placed all over the length of a specimen peak-heated to350◦C at a heating rate of15K/s.2.1.3.Regulation set-upThe experimental device has been designed to reach high heat-ing rates.An accurate control of the temperature is required in order to avoid overshoots.To do so,a PID controller has been used [16–19].The resulting thermal loading is slightly delayed but the heating rate is equal to the desired one.The cooling rate is maxi-mum at the highest temperature(of the order of23K/s at500◦C) and decreases during cooling;it drops to about6K/s when temper-ature becomes lower than150◦C.2.2.Transmission Electron MicroscopyThe experimental device presented previously has been used to heat specimens for both mechanical measurements and TEM observations.Two types of microstructural observations have been carried out during this work.Theﬁrst one is a detailed observation of the microstructure of the material in the T6temper by means of HRTEM(High Resolution Transmission Electron Microscopy) and the second one by means of classical TEM to compare the microstructure of the alloy for three different states of precipita-tion.They were conducted on a JEOL2010F microscope operating at200kV,which belongs to the Centre Lyonnais de Microscopie (CLYM)located at INSA Lyon(France).TEM allows only very local observations so it was not intended to measure accurately the volume fraction of the precipitates;also not enough precipitates were analyzed to obtain an accurate mean radius.2720 D.Maisonnette et al./Materials Science and Engineering A528 (2011) 2718–2724The samples used in TEM are thin lamellas.A disk with a thick-ness of about200␮m is extracted from the heated specimen by means of a diamond wire saw.Its diameter is then reduced by punching.The disk is thinned to electron transparency(thickness to about200nm or less)by electropolishing using an electrolytic bath composed of20%of HNO3in methanol.The bath is cooled at−30◦C with liquid nitrogen[20].A Precision Ion Polishing System(PIPS) is used in order to accomplishﬁnal thinning and cleaning by ion milling.Some EDX(Energy-dispersive X-ray spectroscopy)analy-sis were performed with an Oxford Instruments analyzer,using a nanoprobe(about3nm in diameter)in the TEM to estimate the composition of the precipitates in the T6state.2.3.Mechanical characterizationTensile tests have been carried out at room temperature on spec-imens previously heated to peak temperatures of200,300,400,500 and560◦C with various heating rates(0.5,5,15,50,200K/s)in order to measure their mechanical properties.The thermal loadings are representative of the thermal histories encountered in EB welding.Three parameters have been investigated.Theﬁrst one is the maximum temperature reached during heating at a given heating rate(r=15K/s).The second one is the heating rate for a given max-imum temperature(T=400◦C).The third one is the dwell time at T=560◦C.This last study is not representative of a welding opera-tion but will allow understanding the variation of the mechanical characteristics during holding at a given temperature which cor-responds to the solution treatment of the alloy.For each test,the specimen is heated to the required temperature while compensat-ing for thermal expansion,then it is cooled to room temperature andﬁnally deformed until fracture at a strain rate of10−2s−1.Dur-ing the test,for a strain close to1.5%,an unloading is performed to measure the elastic modulus.3.Results3.1.HRTEM observations of the material in the T6temperThe aim of the HRTEM investigation on the AA6061-T6is to mea-sure the size of some hardening precipitates and to evaluate their composition in order to characterize the microstructure of the ref-erence T6state.The precipitates present in this state are hard to see owing to their very small size and because they are partly coherent with the aluminium matrix.HRTEM is thus mandatory to image the precipitates.Fig.2(a)shows a TEM picture at high magniﬁcation.Two needle-shaped precipitates can be seen:•Theﬁrst one is oriented along the[001]direction.Its cross sec-tion is observed making its diameter measurable accurately.The measure gives a diameter of about4nm.•The second one is oriented along the[100]direction.It is observed lying in the thin foil.The diffractogram,obtained by using Fourier transform,asso-ciated to theﬁrst precipitate is shown in Fig.2(b).In addition to the{200}diffraction spots associated with the aluminium matrix, weak aligned spots prove that the atomic state is partially disor-dered as for pre-␤ phases.At last,an EDX analysis carried out on the needle-shaped precip-itates by means of a3nm probe gives an atomic ratio X Mg/X Si=1.29 (with a standard deviation of0.3).This value is the average result of measurements onﬁve precipitates.3.2.Classical TEM observations of the microstructural changesFollowing the detailed study of the T6temper,the precipitates for various states were observed by means of classical TEM.The aim is to evaluate the evolution of the microstructure(size and vol-ume fraction of precipitates)as a function of the thermal loading previously submitted to the pared to HRTEM,classi-cal TEM is a better way to evaluate the volume fraction because it allows a larger area to be observed at lower magniﬁcation.How-ever classical TEM is worse than HRTEM to measure accurately the diameter of the precipitates because the images at high magniﬁca-tion are often fuzzy(a difﬁculty inherent to the diffraction contrast in conventional TEM).parison of three precipitation statesThe reference microstructure of the T6temper is here compared to states observed after a heating up to300◦C and400◦C at a heat-ing rate of15K/s and no dwell time at the maximum temperature.Fig.3shows three micrographs obtained from representative sample areas for the three investigated states.In the case of the specimen heated to400◦C,some precipitates with a needle shape are present in the picture.These precipitates are very large,with length between65and170nm and a mean value of112nm,and their diameter ranges between5and11nm with a mean value of 7.35nm.The mean values are calculated by taking into account ten precipitates observed on different pictures.However it should be mentioned that the precipitates could be cut by the sample prepa-ration,consequently the length given above should be considered as indicative only.They will be used to compare the precipitation state.In the two other cases,the precipitates are smaller.Their length is between20and40nm with a mean value of29nm for the T6 temper and between15and40nm with a mean value of25nm for the specimen heated to300◦C.Their diameter ranges between 3.75and4.6nm with a mean value of4.45nm for the T6temper and between2and4nm with a mean value of2.6nm for the specimen heated to300◦C.3.2.2.Precipitate volume fraction evaluationThe precipitate size can be measured by means of TEM pictures. However,it is much more difﬁcult to determine the precipitate volume fraction.Indeed,projections obtained by TEM correspond to volumetric observations but the thickness of the sample is not known accurately.In order to get a rough estimate of the precipi-tate volume fraction,TEM micrographs were compared to pictures obtained by modeling.A computer software has thus been devel-oped in Matlab to simulate these images.Based on three simple parameters describing the precipitation state,the program can reproduce a needle-shaped precipitate distribution in a sample with a uniform thickness.The three parameters are the volume fraction(f v),the mean radius of the needle precipitates(r avg)and their mean length(L avg).A Gaussian size distribution is arbitrarily assumed for the radius and the length with a variance of1and36,respectively.The size distributions are discretized in one hundred classes of size.Once the thickness isﬁxed(illustrations will be given here for a100nm thick material),the total volume is calculated and an iterative algo-rithm increases step by step the number of precipitates in each class to obtain the volume corresponding to the desired f v.The pre-cipitates are then shown graphically on a2D view by distributing them uniformly along the three 001 directions of the Al-matrix, which corresponds to the viewing directions of the TEM micro-graphs shown in Figs.2and3.Fig.4compares the precipitation state observed in the specimen heated to300◦C to two modeled states,theﬁrst one with a volume fraction of3%(Fig.4(a))and the second one with a volume fraction of1.6%(Fig.4(c)).It clearlyD.Maisonnette et al./Materials Science and Engineering A528 (2011) 2718–27242721Fig.2.HRTEM observations of needle precipitates in AA6061-T6.(a)Lattice image at high magniﬁcation;(b)diffractogram(numerical Fourier transform)of the micrograph showing diffraction spots(arrows)arising from the precipitate in addition to the square lattice of the aluminium fcc phase along[001].appears that f v=3%is not representative of the real precipitation state because it is too dense.The volume fraction of1.6%is obvi-ously closer to the volume fraction observed by TEM.The same type of study carried out for the two other investigated states givesa similar volume fraction.3.3.Mechanical characterizationAs indicated previously,three parameters have been investi-gated.Theﬁrst one is the maximum temperature reached at a given heating rate(r=15K/s).The second one is the heating rate for a given maximum temperature(T=400◦C).The third one is the dwell time at T=560◦C.3.3.1.Inﬂuence of the maximum temperature reached at constant heating rateTheﬁrst mechanical study carried out at room temperature deals with the inﬂuence of the maximum temperature reached at a given heating rate on the mechanical properties of the AA6061-T6.The maximum temperatures are T=200,300,400,450,500and 560◦C at a heating rate of r=15K/s.The variations of temperature with time for these various thermal loadings are shown in Fig.5. The tensile tests are then conducted at room temperature and the corresponding true stress—logarithmic strain curves are shown in Fig.6.The curves obtained for the heated specimens are compared with the curve obtained for the T6temper without thermal loading (black continuous line).It is found that the thermal loading con-siderably inﬂuences the mechanical properties of the specimens except for a maximum temperature of200◦C for which the curve (not shown in Fig.6)is exactly the same as that of the T6sample. Indeed,the yield stress Rp0.2decreases from278MPa at T=300◦C to 70MPa at T=500◦C.Increasing the temperature further to560◦C, however,does not change the yield stress.Fig.7illustrates this 75%decrease of the yield stress when the maximum temperature is increased from300to500◦C.The measured values are compared to values from the literature[21]for which the maximumtemper-parison of three precipitation states.(a)T6temper;(b)after heating up to300◦C at15K s−1;(c)after heating up to400◦C at15K s−1.All micrographs were taken along a 100 zone-axis of the aluminium matrix.2722 D.Maisonnette et al./Materials Science and Engineering A528 (2011) 2718–2724Fig.4.Modeling of the precipitate distribution for a reached temperature T =300◦C with r avg =2.6nm and L avg =25nm assuming volume fractions of (a)3%and (c)1.6%and comparison with the real precipitate distribution microstructure observed by TEM (b)displayed at the same scale.The volume fraction of 1.6%is obviously closer toreality.Fig.5.Thermal loadings used for the study of the inﬂuence of the reached temper-ature.Fig.6.True stress—logarithmic strain curves for temperatures up to 560◦C.ature has been held during 30min.It shows that the yield stress at ambient temperature is strongly dependent on the peak tempera-ture reached during the thermal loading,without a dwell time at the highest temperature,for peak temperature higher than 200◦C.No data without dwell time at the maximum temperature have been found in the literature.The Young modulus has been also measured for each specimen.It has been measured ﬁrstly at the origin of the stress–strain curve and then during the elastic unloading.A mean value is then calcu-lated.It decreases from 68.7GPa for the T6temper to 65.0GPa for the specimen heated to 560◦C which represents a 5.4%decrease.3.3.2.Inﬂuence of the heating rateThe second mechanical study investigates the inﬂuence of the heating rate on the mechanical properties of the AA6061-T6.The maximum temperature applied here is T =400◦C and the studied heating rates are:r =0.5,5,15,50,and 200K/s.The tempera-ture variation obtained for r =50K/s shows an overshoot of 8◦C which results in a slight decrease of the measured stress.Simi-larly,the temperature of the specimen heated with a heating rate of r =200K/s did not reach T =400◦C but T =362◦C.Consequently,the measured stress for this specimen would be higher than expected.The tensile tests give the true stress—logarithmic strain curves shown in Fig.8.They show that the yield stress Rp 0.2decreasesFig.7.Yield stress variation versus reached temperature from measurements (with-out temperature holding)and from the literature (with a 30min dwell time).D.Maisonnette et al./Materials Science and Engineering A528 (2011) 2718–27242723Fig.8.True stress—logarithmic strain curves for various heating rates up to200K/s. for every heated specimens compared to the T6temper and the lower the heating rate is,the lower the yield stress of the material is.More precisely Rp0.2decreases from170MPa for a heating rate of r=200K/s to96MPa for a heating rate of r=0.5K/s.These values have not been compared with literature since no data dealing with the inﬂuence of the heating rate has been found.3.3.3.Inﬂuence of holding time at560◦CThe last mechanical study accomplished on the material is con-cerned with the inﬂuence of a holding time at high temperature before doing the tensile test at room temperature.This last study compares the mechanical properties of the AA6061-T6after a heat-ing at T=560◦C with and without a dwell time at this temperature. The temperature T=560◦C has been chosen because it is close to the solvus temperature of the␤phase in the␣phase.The chosen dwell time is t=30min and the heating rate is r=15K/s.The mechanical properties obtained for both cases are strictly identical.This result indicates that the dwell time at T=560◦C does not inﬂuence the mechanical properties measured on the tested specimens.4.Discussion4.1.PrecipitationAccording to literature[5–9,13,14],the precipitates which are normally present in the T6temper of the AA6061alloy are very thin and their density is quite high.They are small needles of␤ (or pre-␤ )type.They are oriented following the three 100 matrix directions.Some authors[6,10,22]have carried out a detailed study of the␤ phase.It appears that the X Mg/X Si atomic ratio is very often close to1as reported in Table1.However,other authors[23]man-aged to measure a X Mg/X Si ratio higher than1for GP zones and co-clusters contained in an aged6061.In addition,the observed precipitates are only partially coherent as for the pre-␤ phase. Based on these results,it can be assumed that the precipitates con-tained in the studied reference material are pre-␤ or␤ phases (although the X Mg/X Si atomic ratio measured here to1.29is slightly higher than1).Otherwise,Andersen et al.[6]measured needle pre-cipitates with a size of about4nm×4nm×50nm for the␤ phase and20nm×20nm×500nm for the␤ phase.Furthermore,Don-nadieu et al.[8]measured the size of the precipitates contained in a 6065-T6alloy.They obtained a mean diameter of2.86nm.By com-paring these values to those presented in Sections3.1and3.2it can be assumed that the precipitates contained in the studied AA6061 after heating at400◦C are composed of the␤ phase.On the con-trary,the precipitates contained in the6061-T6and in the6061 after a heating at300◦C are smaller.Therefore,the precipitates are probably remaining␤ precipitates for the6061alloy after heating at300◦C.In addition to that,large intermetallics are visible in the micro-graphs at low magniﬁcation,as shown in Fig.9.The size of the intermetallics ranges from50to300nm.These intermetallics formed during the elaboration of the material do not contribute to the hardening of the alloy.An energy dispersive X-ray spectrometry analysis(EDX)proved that their composition type is(Fe–Cr–Mn–Si) and not(Al–Mg–Si)as for hardening precipitates.The structure of these intermetallics was not investigated further.However,it is important to note that the intermetallics do contain silicon,so that the corresponding quantity will not be available for hardening precipitation.4.2.Mechanical propertiesFig.6showed that the behavior of the material after heating at500◦C is strictly identical to the behavior of the material after heating at560◦C.Thus,it can be assumed that the microstructure is the same in both cases.Furthermore,a tensile test carried out on a specimen heated to560◦C during30min gives exactly the same behavior.This behavior corresponds to the O temper.It is commonly accepted that a long holding time at T=560◦C(solvus temperature of the␤phase in the␣phase)is required to dis-solve the parison of the true stress—logarithmic strain curves obtained with and without dwell time shows that the mechanical properties are identical.This means that the dwell time at T=560◦C does not change the mechanical properties.The microstructure is therefore identical corresponding to the annealed state(or O temper)for which no precipitate is present in the mate-rial.This last result shows that for the heating rate and for the specimens used in this study,it is not necessary to apply a dwell time to reach the O temper.This conclusion is probably not valid in the case of a large structure since the peak temperature at each point within the material would depend on its distance from the closest surface.Another result of this investigation is that the heat-ing rate has an inﬂuence on the mechanical properties.By using a higher heating rate,the O temper could not be obtained without a dwell time.The hardening is due to the precipitates contained within the material.They hinder dislocation glide.For a given volume fraction, hardening is most effective if the precipitates are small(and there-fore more numerous).These small precipitates have been observed by TEM for the T6temper.This microstructure leads to more favor-able mechanical properties than the other investigated states.The behavior observed here is quite close to the one observed by Zain-ul-Abdein et al.[24]on a6056-T4.Then,the microstructure of the specimen heated to300◦C seems to be close to the one observed for the T6temper,which explains the small difference of mechanical properties.If the maximum temperature is further increased,the yield stress Rp0.2decreases signiﬁcantly as shown in Fig.7.The TEM observations show that this decrease is due to a strongly enhanced growth of the precipitates.The volume fraction of the precipitates remains identical so that the precipitate number is decreased.This results in a sharp decrease of the mechanical properties,as high-lighted by the tensile tests.Concerning the study of the inﬂuence of the heating rate,no microstructural observations have been carried out.However,Fig.8 shows a decrease of the mechanical properties for every thermal loading up to400◦C compared to the mechanical properties of the T6temper.This means that the material has encountered a microstructural change for every investigated heating rate.If the heating rate is very low,the microstructural changes as dissolution and growth of precipitates,have more time to occur.Consequently, less precipitates are present(for a constant volume fraction)and the mechanical properties are lower.The Young modulus has been measured and it has been shown that it decreases slightly compared to the T6state when。

汽车结构用590MPa级高屈服强度钢的研制_王科强

study on corrosion protection ［D］． Hefei： Hefei University of ［23］Palomino L M，Suegama P H，Aoki I V，et al． Electrochemical study
Technology，2009： 15-53．
alloys，and related materials for biomedical applications［J］． Materials
［20］张林，谭帼馨，宁成云，等．钛表面自组装硅烷化对仿生矿化
Science and Engineering Ｒ，2004，47（ 3）： 49-121．
表 1 汽车结构用 590 MPa 级高屈服强度钢的化学成分（质量分数，%）
Table 1 Chemical composition of 590 MPa grade high yield strength steel for automobile structure （ mass fraction，%）
连铸坯在步进式加热炉中加热，为增加 Nb 和 Ti 的固溶量，提高析出强化效果，加热温度 ≥1250 ℃ 。粗轧后高压除鳞在 1780 机组上精轧成 3 ～ 5 mm 厚板卷。终轧温度≥890 ℃ ，轧后前段层流冷却，卷取温度 ≥650 ℃ 。热轧原料经酸洗后，在冷轧联合机组进行轧制，此使钢组织中的珠光体团间距减小和珠光体被破碎得较充分，为退火过程中的晶粒细化和再结晶提供条件，冷轧压下率在 55% ～ 70% 之间。 2． 3 连续退火
and osteoblast proliferation ［ J ］．Ｒare Metal Materials and
automotive steels［J］． Surface Engineering，2000，16（ 4）： 315-320．

原位拉伸研究热处理对激光选区熔化GH4169_合金组织及650_℃_力学性能的影响

2024 年第 44 卷航　空　材　料　学　报2024，Vol. 44第 1 期第 93 – 103 页JOURNAL OF AERONAUTICAL MATERIALS No.1 pp.93 – 103引用格式：朱嘉冕，吕国森，姜文祥，等. 原位拉伸研究热处理对激光选区熔化GH4169合金组织及650 ℃力学性能的影响[J]. 航空材料学报，2024，44(1)：93-103.ZHU Jiamian，LYU Guosen，JIANG Wenxiang，et al. Effect of heat treatment on microstructure and mechanical properties of selective laser melting GH4169 alloy at 650 ℃: in-situ SEM investigation[J]. Journal of Aeronautical Materials，2024，44(1)：93-103.原位拉伸研究热处理对激光选区熔化GH4169合金组织及650 ℃力学性能的影响朱嘉冕1, 吕国森1, 姜文祥1, 程晓鹏1, 贾泽一1, 吕俊霞1*,黄帅2, 张学军2（1．北京工业大学材料与制造学部，北京 100124；2．中国航发北京航空材料研究院，北京 100095）摘要：研究热处理制度对激光选区熔化成形GH4169合金组织及高温力学性能的影响。

通过自主研发的SEM原位加热拉伸测试平台，探究热处理前后650 ℃合金力学性能变化与动态组织演变的关系。

结果表明：热处理后合金的晶粒形态由柱状晶转化为等轴晶，Laves相溶解，析出大量γ′和γ′′强化相；在650 ℃下，沉积态合金的屈服强度和抗拉强度分别为574 MPa和740 MPa，热处理态（HSA态）合金的屈服强度和抗拉强度分别为818 MPa和892 MPa，较沉积态分别提升了42.5%和20.1%；沉积态合金表面晶粒起伏更大，协调变形能力更强，塑性流动能力好；裂纹在Laves相周围萌生沿枝状晶向最大切应力方向扩展，样品颈缩后发生剪切断裂；HSA态裂纹在碳化物周围萌生沿晶界扩展，断裂方式为沿晶和穿晶相结合的混合断裂。

等温锻造钛合金技术研究新进展

等温锻造钛合金技术研究新进展罗谨灵;杨启明;毛东壁;冯斐斐【摘要】钛合金强度高、密度低、耐热耐腐蚀、韧性好,是良好的现代工业和航空航天材料.采用等温锻造工艺的钛合金变形均匀,微观结构较好,加工余量小,机械性能高于常规锻造.文中介绍了近10年等温锻造钛合金材料研究及产品开发进步现状,总结了温度和应变速率在钛合金等温锻造中的影响效果.同时结合新材料和新工艺分析了今后钛合金等温锻造研究的发展方向.【期刊名称】《机械工程师》【年(卷),期】2015(000)001【总页数】3页(P17-19)【关键词】等温锻造;钛合金;应变速率;温度【作者】罗谨灵;杨启明;毛东壁;冯斐斐【作者单位】西南石油大学,成都610500;西南石油大学,成都610500;西南石油大学,成都610500;西南石油大学,成都610500【正文语种】中文【中图分类】TG3160 引言“等温锻造”的概念出现于20世纪60年代，其研究始于美国，很快前苏联也投入其中，使得相关工作取得进步。

等温锻造是指在锻造过程中，毛坯和模具都保持稳定或变化非常缓慢且高于常规锻造的温度，创造出优于常规锻造的条件，使许多常规锻造存在的缺陷在等温锻造条件下能较好地克服。

钛是20世纪50年代开发出供制造业使用的金属，具有强度高、耐热耐腐蚀、韧性好等优点，是汽轮机、发动机及航空航天等精密零部件的优选材料。

据NASA估计，到2020年发动机材料总量的20%～25%将是钛铝合金。

钛合金锻造融合CAD/CAM技术，采用等温锻造、超塑性锻造等新兴工艺，可防止出现在常规锻造条件温度下降迅速导致材料成形困难的状况，使毛坯保持和模具相同的温度，提升钛合金材料塑性，可将需要多次成形的工件一次精锻成形。

除此之外，等温锻造减少了锻后加工余量，节省购买昂贵的钛合金材料而支付的生产成本，同时降低锻造效果对操作人员个人技术的依赖，大幅提升锻件质量。

1 等温锻造特点常规锻造下，毛坯温度高于模具温度并不断将热量传递给模具，毛坯自身迅速冷却，内部分子动能降低，难以克服分子间引力，导致塑性降低，变形抗力增加。

热冲压用钢22MnB5回火组织与性能分析

ＷｕＧｕａｎｈｕａ１ꎬ２ ꎬＬｉＪｉａｎｙｉｎｇ２ ꎬＭａＧｕａｎｇｚｏｎｇ２ ꎬＳｕｎＬｕ２ ꎬＳｕｎＨｏｎｇｌｉａｎｇ２
(１. ＣｏｌｌｅｇｅｏｆＭａｔｅｒｉａｌｓＳｃｉｅｎｃｅａｎｄＥｎｇｉｎｅｅｒｉｎｇｏｆＮｏｒｔｈＣｈｉｎａＵｎｉｖｅｒｓｉｔｙｏｆＳｃｉｅｎｃｅａｎｄＴｅｃｈｎｏｌｏｇｙꎬ Ｔａｎｇｓｈａｎꎬ Ｈｅｂｅｉꎬ ０６３００９ꎻ ２. ＴｅｃｈｎｉｃａｌＣｅｎｔｅｒｏｆＨＢＩＳＧｒｏｕｐＴａｎｇｓｔｅｅｌＣｏｍｐａｎｙ ꎬ Ｔａｎｇｓｈａｎꎬ Ｈｅｂｅｉ ꎬ ０６３０１６) Ａｂｓｔｒａｃｔ: Ｖｅｈｉｃｌｅｌｉｇｈｔｗｅｉｇｈｔｉｓｏｎｅｏｆｔｈｅｋｅｙｔｏｐｉｃｓｉｎｔｈｅａｕｔｏｍｏｔｉｖｅｉｎｄｕｓｔｒｙ. Ｔｈｅｕｓｅｏｆｕｌｔｒａｈｉｇｈｓｔｒｅｎｇｔｈｓｔｅｅｌｃａｎｅｆｆｅｃｔｉｖｅｌｙｒｅｄｕｃｅｔｈｅｂｏｄｙｗｅｉｇｈｔ. Ｉｎｔｈｉｓｐａｐｅｒꎬ ｔｈｅｅｆｆｅｃｔｏｆｔｅｍｐｅｒｉｎｇｈｅａｔｔｒｅａｔｍｅｎｔｏｎｍｉｃｒｏｓｔｒｕｃｔｕｒｅａｎｄｍｅｃｈａｎｉｃａｌｐｒｏｐｅｒｔｉｅｓｏｆ２２ＭｎＢ５ｓｔｅｅｌｐｒｏｄｕｃｅｄｂｙＴａｎｇｓｔｅｅｌｉｓｓｔｕｄｉｅｄ. Ｔｈｅｒｅｓｕｌｔｓｓｈｏｗｔｈａｔｔｈｅｔｅｍｐｅｒｉｎｇｈｅａｔｔｒｅａｔｍｅｎｔｃａｎｅｆｆｅｃｔｉｖｅｌｙｒｅｄｕｃｅｔｈｅｉｎｔｅｒｎａｌｓｔｒｅｓｓｏｆｔｈｅ２２ＭｎＢ５ａｆｔｅｒｑｕｅｎｃｈｉｎｇｔｒｅａｔｍｅｎｔꎬ ｐｒｅｖｅｎｔｔｈｅｂｒｉｔｔｌｅｆａｉｌｕｒｅꎬ ａｎｄｔｈｅｌｏｗｔｅｍｐｅｒａｔｕｒｅｔｅｍｐｅｒｉｎｇｃａｎｅｆｆｅｃｔｉｖｅｌｙｉｍｐｒｏｖｅｔｈｅｃｏｍｐｒｅ￣ｈｅｎｓｉｖｅｐｒｏｐｅｒｔｉｅｓｏｆｔｈｅｍａｔｅｒｉａｌ. ＫｅｙＷｏｒｄｓ: ｈｏｔｓｔａｍｐｉｎｇｓｔｅｅｌꎻ ｔｅｍｐｅｒｉｎｇꎻ ｍｅｔａｌｌｏｇｒａｐｈｉｃｓｔｒｕｃｔｕｒｅꎻ ｐｈｙｓｉｃａｌｐｒｏｐｅｒｔｉｅｓ

调质热处理对EH47船板钢显微组织的影响

第43卷第1期722021年1月上海金属SHANGHAI METALSVol.43,No.1January，2020调质热处理对EH47船板钢显微组织的影响李洪楠1张红梅“赵大东1王渐灵2刘焕然1李娜1姜正义2 (•辽宁科技大学材料与与金学院，辽宁鞍山114051；2.鞍钢集团朝阳钢铁有限公司，辽宁朝阳122000;3.海洋装备用金系材料及其应用国家重点实验室，辽宁鞍山114009)【摘要】在实验室条件下研究了层质工艺参数(淬火温度870~960C，回火温度580~ 670C)对EH47船板钢显微组织的影响。

结果表明：试验钢经经质处理后的显微组织是以针状铁素体为主，含一定量粒状铁素体、准多边形铁素体和粒状贝氏体的混合组织。

回火温度相同,随着淬火温度的升高，试验钢的组织更为细小均匀，贝氏体含量稍有增加;淬火温度相同，随着回火温度的升高，组织更为细小，针状铁素体含量增加，贝氏体含量减少。

930C保温30min 淬火、640C回火30min的钢的组织最为细小均匀。

因此,EH47船板钢的实际际质淬火温度应高于900C，回火温度控制在580~640C。

【关键词】EH47船板钢调质处理针状铁素体贝氏体Effect of Quenching and Tempering Processes onMicrostructure of EH47Hull SteelLI Hongnan1ZHANG Hongmei1，ZHAO Dadong1WANG Jianling2LIU Huanran1LI Na1JIANG Zhengyi1，(1.School of Materials and Metallurgy,University of Science and Technology Liaoning,Anshan Liaoning 114051,China;2.Anshan Steel Group Chaoyang Steel Company Limited,Chaoyang Liaoning122000, China;3.State Key Laboratory of Metal Material for Marine Equipment and Application,Anshan Liaoning114009，China)[Abstract]Effect of quanching-and-tempering process parameters,that is,austenitizing temperatures of870to960C and tempering temperatures of580to670C,on microstructure of EH47hull steel was investigated in laboratory.The results showed that microstructure of the steel after quenching and tempering was predominantly acicular ferrite,also contained a certain amount of granular ferrite,quasi-polygonal ferrite and granular bainite.With the increase in austenitizing temperature,the steel exhibited even finer and even more uniform microstructure,and slightly decreased bainite content after tempering at the same temperature.With the increase in tempering temperature,the steel quenched from the same temperature exhibited even finer microstructure which contained more acicular ferrite and less bainite.The steel held at930C for30min and quenched, then tempered at640C for30min given the finest and most uniform microstructure.In view of the foregoing,actual austenitizing temperature should be above900C,and tempering temperature should be controlled in the range of580to640C for the EH47hull steel.【Key Words]EH47hull steel,quenching-and-tempering,acicular ferrite,bainite基金项目：辽宁省自然科学基金(20180550952);辽宁科技大学与海工钢国家重点实验室联合项目(SKLMEA-USTL2017010和201905)作者简介:李洪楠，女，主要从事钢铁材料组织性能控制研究,E-mail：lihongnan7270@,电话=158****7270通信作者:张红梅，女，教授，博士，主要从事钢铁材料组织性能控制和微成形理论与工艺研究,E-mail：lilyzhm68@第1期李洪楠等:调质热处理对EH47船板钢显微组织的影响73近年来，由于海洋贸易日益频繁，海洋工程建设与深海探索活动蓬勃发展，船舶更新换代周期越来越短，世界市场对船板钢的需求与日俱增［-3］o船体结构用钢是船舶制造的重要材料，主要用于制造远洋及内河航运船舶的船体、甲板等。

影响烧结Nd-Fe-B磁体退磁曲线方形度的因素

影响烧结Nd-Fe-B磁体退磁曲线方形度的因素王占勇1，谷南驹1，王宝奇1，刘金芳2，赵金伶3，张志清3，张巧格3（1．河北工业大学金属材料研究所，天津300132；2．美国宾夕法尼亚洲电子能公司，宾夕法尼亚州17538，美国；3．河北省冶金科技股份有限公司磁材部，河北石家庄050000）摘要：通过分析具有不同退磁曲线方形度的磁体发现，烧结体的显微组织对磁体的方形度有很大影响。

磁体中晶粒的异常长大会严重恶化磁体的方形度；晶粒的形状及晶界相等影响到退磁场的大小，进而影响到磁体的方形度；添加元素影响到磁体中的相结构和相分布，对反磁化场的均匀性有所影响。

关键词：Nd-Fe-B磁体；方形度；晶粒；显微组织；添加元素1引言Nd-Fe-B是当代磁能积最高的永磁材料，被称为“磁王”。

目前，对这种高性能磁体的研究主要朝两个方向进行，一是高磁能积磁体，日本实验室水平已达444kJ/m3，工业批量生产水平为N50[1]（磁能积400kJ/m3）；一是高矫顽力和低温度系数磁体，这一类磁体主要用在电机等领域，前景很好。

然而，在实际应用中，仅仅考虑磁能积和矫顽力这两个指标是不够的，还必须考察磁体的退磁曲线方形度（以下简称方形度）是否合乎要求。

图1为典型的永磁体的退磁曲线[2]，从J～H曲线上我们看出，在反向（退）磁场比较小时，J的下降很小；反向磁场大到一定程度后，J开始急剧下降。

通常把J=0.9B r或0.8B r的退磁场称为弯曲点磁场H k。

H k/H cj在一定程度上反映了J～H退磁曲线的形状，其比值越接近于1，J～H退磁曲线越接近于方形，所以，生产中经常通过比较H k/H cj的大小来衡量方形度的好坏，这种衡量方法在许多文献[3，4]中都被采用。

通常认为方形度HH cj＞0.9，产品就算合格。

k/在生产中经常发现方形度不合格的产品，我们对这些情况出现的原因进行了分析，总结出了影响方形度的一些因素，以供大家参考。

本文中涉及到的H k 都是指J=0.9B r所对应的磁场。

法国N19_新型粉末高温合金的研发动态及进展

件寿命显著增加。这项技术最初由美国开发，大多镍基高温为增加 γ 基体的固溶强化作用，适当降了 Mo 并添加 W。通
合金都是为某型发动机专门设计的。法国拥有独立自主的高过适量提高 Cr 的含量来增强高温抗氧化和耐腐蚀性，避免
温合金研发路线，主要是由斯奈克玛公司（Snecma）主导推动，并成功将新材料应用于飓风战斗机的 M88 发动机中 [1]。法国粉末高温合金的研发具有自己独特的思路和方式，例如 N18、NR3、N16 等盘合金的成功设计和应用标志着法国在国际高温新材料的研发中具有一席之地。目前，N19 合金是法国最新的粉末高温合金，本文将综述该合金的设计、工艺、显微组织与性能，为国内粉末高温合金设计与开发人员提供借鉴。
表 3 Rene88DT、N18 和 N19 粉末高温合金蠕变性能对比
合金牌号 Rene88DT
N18
N19
蠕变条件 700/550 700/700 700/550 650/~650 700/700 700/650 650/~925
蠕变应变到达0.2% 的时间/h 340 260 585 ~30 380 462 ~30
A CD
B E
（a）曲线意义示意图
工业技术
（b）第一种热处理 γ' 相形貌
（c）第二种热处理 γ' 相形貌
（d）第三种热处理 γ' 相形貌时间/min （e）第四种热处理 γ' 相形貌
- 84 -
（f）第五种热处理 γ' 相形貌图 4 不同热处理工艺方法对 N19 合金 γ' 相形貌和尺寸的影响
图 2 经过 HIP+HEX+ITF 后的 N19 合金变形显微组织特征
- 82 -

219515817_带状组织对微合金钢应变局部化及应变强化行为的影响

精密成形工程第15卷第6期120 JOURNAL OF NETSHAPE FORMING ENGINEERING2023年6月收稿日期：2023–02–05 Received ：2023-02-05 基金项目：天津市自然科学基金（JCQNJC03700）；国家自然科学基金（11872275）；大学生创新创业训练项目（202110069075） Fund ：The Natural Science Foundation of Tianjin(JCQNJC03700); The National Natural Science Foundation of China (11872275); Innovation and Entrepreneurship Training Program for College Students(202110069075) 作者简介：任春华（1989—），男，博士，讲师，主要研究方向为实验固体力学。

Biography ：REN Chun-hua(1989-), Male, Doctor, Lecturer, Research focus: experimental solid mechanics. 引文格式：任春华, 周治顺, 薛冬阳, 等. 带状组织对微合金钢应变局部化及应变强化行为的影响[J]. 精密成形工程, 2023, 15(6): 120-126.REN Chun-hua, ZHOU Zhi-shun, XUE Dong-yang, et al. Effect of Banded Microstructure on Strain Localization and Strain 带状组织对微合金钢应变局部化及应变强化行为的影响任春华，周治顺，薛冬阳，张晓川，计宏伟，寇金宝（天津商业大学机械工程学院，天津 300134）摘要：目的研究不同加载方向下带状结构对微合金钢塑性变形及应变强化行为的影响。

镍基高温合金长期时效过程中第二相的析出

3
结
论
（l）合金在长期时效时会析出碳化物相与针状 TCP 相・碳化物与 TCP 相的析出特性有着类似之处：其颗粒尺寸大于 800C 时效容易析出， 900C 时效时的尺寸・（2）实验观察到析出碳化物种类有 M 23 C6 型、 M 7 C3 型； TCP 相有 " 相、 # 相两种・（3）趋向 !' 相的形貌随着时效时间的延长，于方形，而分布则越来越不平衡，聚集并长大・基体中析出的 !' 相的粒子平均半径的负三次方与时效时间成正比，符合里夫希茨・瓦格纳的第二相粒子成熟理论・参考文献：
能的显著下降，是合金组织中所不希望出现
［ 7， 8］的本实验主要通过对固溶处理后的试样进・
行不同的温度与时间的时效处理后研究合金在不同长期时效制度下的 #' 相、碳化物和 TCP 相的析出规律・
1
实验方法
实验用料为高温合金板材，取自抚顺钢厂提
再在 M TP-1 磁力驱动双喷电 mm 小圆片各 5 个，解减薄器上最终减薄，电解液为 10% HCIO4 + 抛光电压 50 ~ 75 V ，电流 50 mA， 90% C2 H5 OH，双喷时利用干冰将电解温度控制在 - 25 ~ - 30 利用透射电镜进行观察并照相、确定晶体 C 之内・结构・ !"# 扫描电镜观察利用线切割方法切下大小适中的所需试样・在由粗到细的砂纸上仔细磨光并抛光・利用腐蚀液：（20 mI）+ C2 H5 OH （ 80 mI）+ CuSO（ HCI 4 4 g）进行腐蚀，直至在光学显微镜下看到没有明显的划痕，并且可以清晰地看到晶界与部分晶内析出相（长期时效试样）为止，利用扫描电镜进行观察、电子探针分析并照相・

Zr_Sn_Nb_Fe合金金属间化合物及其_相变温度的研究

第30卷第1期2009年 2月材料热处理学报TRANSACTIONS OF MATERIALS AND HEAT TREATMENTVol .30 No .1February 2009Zr Sn Nb Fe 合金金属间化合物及其相变温度的研究梁建烈1,2, 唐轶媛1, 严嘉琳2, 朱其明1, 庄应烘2, 文国富1(1 广西民族大学物理与电子工程学院,广西南宁 530006;2 广西大学材料科学研究所,广西南宁 530004)摘要:利用大功率X 射线衍射仪研究了Zr 1 0Sn 1 0Nb 0 3Fe 合金金属间化合物的晶体结构,结果表明,Zr 1 0Sn 1 0Nb 0 3Fe 合金的金属间化合物是T i 2Ni 结构的立方(ZrNb)2Fe 和六方Zr(NbFe)2。

在轧制状态下,六方Zr (NbFe)2是主要的析出相。

400 淬火时,立方(ZrNb)2Fe 是主要的析出物。

750 时,不再观察到六方Zr(NbFe)2;850 时,不再观察到立方(ZrNb)2Fe 。

综合DTA 和X 射线测试结果,确定合金 Zr Zr 转变开始温度为633 ,这个温度也是Zr (NbFe)2开始熔入 Zr 的温度,在785 ,立方(ZrNb)2Fe 开始熔入 Zr,到1000 时, Zr Zr 转变结束。

关键词:锆合金; 相变; 中间相; 锆; 铌; 铁中图分类号:TG146 4 文献标识码:A 文章编号:1009 6264(2009)01 0032 04Investigation of intermediate phases and phase transitiontemperature of in the Zr Sn Nb Fe alloyLIANG Jian lie 1,2, TANG Yi yuan 1, YAN Jia lin 2, Z HU Qi ming 1, ZHUANG Ying hong 2, WE N Guo fu 1(1 Sc hool of Physics and Electronic Engineering,Guangxi University for Nationalities,Nanning 530006,China;2 Institute of Materials Sciences,Guangxi University,Nanning 530004,China)Abstract :Intermetallics of the Zr 1 0Sn 1 0Nb 0 3Fe alloy was identi fied to be cubic T i 2Ni struture (ZrNb)2Fe and hexagonal Zr(NbFe)2by X ray diffraction.For the as received alloy,the hexagonal Zr(NbFe)2is the major intermediate phase.For the alloy annealed at the temperature higher than 400 ,cubic (ZrNb)2Fe is the main secondary phase particle.No Zr(NbFe)2phase is detected at temperature hi gher than 750 ,and the di ffraction peaks of the (ZrNb)2Fe disappears when the temperature is higher than 850 .Combining DTA and XRD results,the Zr Zr transformation occurs at 633 accompanying the dissolve of the Zr(Nb Fe)2,and wi th temperature increasing,(ZrNb)2Fe starts to dissolve into Zr at 785 ,and the Zr Zr transformation finishs at about 1000 .Key words :zirconium based alloys;transformation;intermetallics;Zr;Nb;Fe收稿日期: 2007 11 14; 修订日期: 2008 11 11基金项目:核燃料及材料国家级重点实验室基金(00J S85 9 1GX0101);广西科学基金(0448022、0728060)以及广西大型仪器协作共用网联合资助作者简介: 梁建烈(1971!),男,广西民族大学材料科学研究所副教授,硕士,发表论文10多篇,主要从事合金相图与相变研究,电话:0771 *******,E mail:li angjl1971@126 c om 。

高导磁坡莫合金一大缺点就是磁性对应力极为敏感

2564 0.097
78%
404 Hv0.5 138 Hv0.3 150 Hv0.3
2564 0.114
76%
407 Hv0.5 136 Hv0.3 146 Hv0.3
2565 0.114
76%
403 Hv0.5 141 Hv0.3 156 Hv0.3
2565 0.143
68%
396 Hv0.5 135 Hv0.3 146 Hv0.3
PC-47 合金的研制
牛永吉 1, 2，李振瑞 1，陈国钧 1 （1. 北京北冶功能材料有限公司 100192；2. 北京科技大学）
摘要：研究了双声道、四声道音频磁头芯片用 Ni81Nb3.8Mo1.1Fe（牌号 PC-47）坡莫合金的成分和工艺，考察了热处理工艺、树脂固化对磁性能的影响。研究说明，所研制的合金同进口料性能相当，可完全替代进口材料。研究说明，该料严格的成分控制与工艺控制，是确保材料性能稳定一致的关键，该料成分控制的关键在于 Ni 含量及合金元素特别是 Nb 含量的稳定控制，合金最佳成分为 Ni80.8Nb3.8Mo1.1Fe；材料软态硬度随合金元素特别是 Nb 含量增加而增加；应力对磁性的影响程度与 Ni、Nb 含量的关系较大；该料直流性能随退火温度升高而提高，交流磁性能则有一最佳退火温度，综合考虑该料最佳退火温度在 1100~1150℃；该料在保护气氛中退火后在大气环境中的时效氧化处理有利于提高硬度，改善交流性能，改善磁性应力敏感性，适于音频磁头应用。本合金的开发在生产适用于多声道音频磁头的高性能稳定性、低磁性应力敏感性、高精度软磁材料上实现了突破，打破了国外垄断。实际使用表明该材料是一种有竞争力的廉价新型磁头材料，另外，该料在受应力作用下使用软磁性能的环境中值得推荐。

Zr-Sn-Nb-Fe-V合金包壳管加工过程中第二相的演变

Zr-Sn-Nb-Fe-V合金包壳管加工过程中第二相的演变吴宗佩;易伟;杨忠波;程竹青;陈波全【摘要】采用扫描电子显微镜(SEM)和透射电子显微镜(TEM)研究了Zr-0.2Sn-1.3Nb-0.2Fe-0.05V合金经热挤压、冷轧、中间退火包壳管坯以及经终轧及最终退火后成品管材第二相特征.结果表明,热挤压产生的β-Zr及第二相沿管坯轴向呈流线状分布,随着冷轧和退火的进行,亚稳相β-Zr发生分解,第二相逐渐均匀化,最终呈细小、均匀、弥散分布.合金成品管材第二相主要为BCC结构的β-Nb,含有少量FCC结构的Zr(NbFeV)2.加工过程中析出相的平均直径变化不大,均小于100 nm.合金包壳管第二相尺寸分布与热处理过程中含Nb第二相溶解析出直接相关.【期刊名称】《原子能科学技术》【年(卷),期】2019(053)003【总页数】7页(P385-391)【关键词】Zr-Sn-Nb-Fe-V合金;包壳管;加工过程;第二相【作者】吴宗佩;易伟;杨忠波;程竹青;陈波全【作者单位】中国核动力研究设计院反应堆燃料及材料重点实验室,四川成都610213;中国核动力研究设计院反应堆燃料及材料重点实验室,四川成都 610213;中国核动力研究设计院反应堆燃料及材料重点实验室,四川成都 610213;中国核动力研究设计院反应堆燃料及材料重点实验室,四川成都 610213;中国核动力研究设计院反应堆燃料及材料重点实验室,四川成都 610213【正文语种】中文【中图分类】TL341锆合金因具有低的热中子吸收截面、良好的耐腐蚀性能、适中的力学性能、优良的耐辐照性能[1]，是目前压水堆核电站唯一使用的包壳材料。

随着压水堆朝高燃耗、长换料周期方向的发展，现有包壳材料Zr-4合金已无法满足高燃耗燃料组件的要求。

为此，世界各国都开展了高性能新锆合金的研制，其中Zr-Sn-Nb-Fe合金是压水堆高燃耗组件用锆合金持续改进的重要方向。