Electrochimica Acta 205 (2016) 178–187

Thermal expansions and mechanical properties of electrodepositedFe–Ni alloys in the Invar composition range

Tomio Nagayama*, Takayo Yamamoto, Toshihiro NakamuraKyoto Municipal Institute of Industrial Technology and Culture, 91 Chudoji Awata-cho, Shimogyo-ku, Kyoto 600-8815, Japan

A R T I C L E I N F O

Article history:Received 10 January 2016Received in revised form 15 April 2016Accepted 18 April 2016Available online 22 April 2016

Keywords:Electrodeposition of Fe-Ni alloyInvar alloyThermal expansionMechanical propertySulfur embrittlement

A B S T R A C T

Electrodeposited Invar Fe–Ni alloys with 36 to 40 mass% Ni were prepared from plating baths containingsaccharin as a stress reducer and containing various Fe2+ concentrations. The Invar Fe–Ni alloys containedof small amount of S (�0.02 mass%). The coefficients of thermal expansion (CTEs) of the as-depositedInvar Fe–Ni alloys were approximately 9 to 11 ppm/�C and were larger than those of pyrometallurgicallyproduced Invar alloys. When the alloys were heat-treated at 400 to 500 �C, their CTEs drasticallydecreased to approximately 5 ppm/�C. Furthermore, upon heat treatment at 600 �C, the CTEs reachedapproximately 2 to 4 ppm/�C depending on alloy composition; these CTEs are comparable with those ofpyrometallurgically produced alloys. The as-deposited Invar Fe–Ni alloys were mainly composed ofmetastable bcc phases, resulting in larger CTEs. When the alloys were annealed at 400 �C or above, theequilibrium fcc phases became the predominant phases, accompanied by a drastic decrease of the CTEs.The bcc-to-fcc transformation led to a decrease of the CTEs and to thermal contractions. Upon the heattreatment, an S (sulfide) at bcc grain boundaries segregated not as a thin film but as a granular sulfideform at primary bcc grain boundaries in the electrodeposited Invar alloys. Upon heat treatment at 500 �Cor above, bcc grain eliminated accompanied by fcc grain growth and the granular sulfide agglutinatedfurther. Consequently, the agglutinating granular sulfide was entrapped in the matrix grains or grain-boundary triple-points of transformed fcc grains. In addition, it is considered that these two-typemorphologies of the sulfide, as a thin film at grain boundaries or as a precipitate, will be determined bygrain growth form during an annealing. Ductile behavior of the electrodeposited Invar Fe–Ni alloys wasconfirmed, irrespective of whether the alloys were heat-treated. Upon heat treatment at 400 to 500 �C,the Invar Fe–Ni alloys exhibited high strength with good ductility, consistent with their low CTE. After theheat treatment, no severe embrittlement of the electrodeposited Invar alloys was observed despite thecodeposition of S because the S existed as a granular sulfide, thereby preventing grain-boundaryembrittlement.

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1. Introduction

At temperatures near room temperature, Fe–Ni alloys with30 to 40 mass% Ni have coefficients of thermal expansion (CTEs)lower than those of pure Fe and Ni (12 and 13 ppm/�C, respectively)[1,2]. In particular, Fe-36 mass%Ni alloy exhibits the lowest CTE inthe Fe–Ni alloy system (approximately 1 ppm/�C); this is known asthe Invar effect, which was discovered by Guillaume [1] in 1897;Fe-36 mass%Ni alloy and Fe–Ni alloys of similar composition arethus referred to as “Invar alloys.” The Invar alloys are used innumerous applications such as optical and laser measuring

* Corresponding author.E-mail address: nagayama@tc-kyoto.or.jp (T. Nagayama).

http://dx.doi.org/10.1016/j.electacta.2016.04.0890013-4686/ã 2016 Elsevier Ltd. All rights reserved.

systems, bimetallic strips, shadow masks, and storage vesselsfor liquefied natural gas among others [2].

Electrodeposition of the Invar Fe–Ni alloys is expected to enablemore precise processing with higher throughput and result inalloys with improved mechanical properties compared with thoseof Invar alloys prepared by conventional processes such asrolling, machining, and etching methods. In addition, Invarelectroforming processes can provide freestanding micrometer-sized, 3-dimensional structures with good thermal dimensionstabilities, e.g., MEMS (microelectromechanical systems) [3–6].Therefore, researchers have focused extensively on electrodepos-ited Invar alloys [3–12,15].

Thus far, electrodepositions of Fe–Ni alloys with Invar(�36mass%Ni [7–10]) and Invar-like compositions (e.g., 42mass%Ni [11]) have been demonstrated. In these previous studies,

Table 1Bath composition and plating conditions [10].

NiSO4�6H2O 0.95 mol/LNiCl2�6H2O 0.17 mol/LBoric acid 0.49 mol/LFeSO4�7H2O 0–0.35 mol/LSaccharin sodium 0.008 mol/LMalonic acid 0.05 mol/LBath pH 2.3Bath Temperature 50 �CCurrent density 40 mA/cm2

T. Nagayama et al. / Electrochimica Acta 205 (2016) 178–187 179

saccharin was added to the plating baths as a stress reducer.Saccharin has also been reported to be indispensable in theproduction of electrodeposited Fe-rich Fe–Ni alloys, e.g., the Invaralloys with �36 mass% Ni. Electrodeposited Invar Fe–Ni alloysdeposited in the absence of saccharin are burnt, torn, and crackedbecause of very high internal stresses [8,12]. We have alreadysuccessfully produced a crack-free Invar Fe–Ni alloy with 36 mass%Ni and have evaluated its thermal properties and microstructure[10].

As-deposited Invar alloys have been reported to be composed ofmetastable body-centered cubic (bcc) phases or mixed bcc andface-centered cubic (fcc) phases; this composition differs fromthose of pyrometallurgically produced Invar alloys, which arecomposed primarily of equilibrium fcc phases [7–10,13,14].Consequently, as-deposited Invar alloys exhibit CTEs larger thanthose of pyrometallurgically produced Invar alloys [7,9,10]. Weused an electroforming process to prepare a fine-pitch Invar metalmask for large- and fine-pitch organic light-emitting diodedisplays [6]. Annealing at 600 �C decreased the CTE of theelectroformed Invar metal mask to 3 ppm/�C, which is four timessmaller than that of a conventional Ni mask (13 ppm/�C) [6]. Inaddition, the phase of the electrodeposited Invar Fe–Ni alloys isknown to be dependent on the experimental conditions, e.g.,current density, current wave form, and additive concentration[7,8]. Consequently, a heat treatment at 400 �C or above isnecessary to decrease the CTEs of the electrodeposited Invaralloys to those of pyrometallurgically produced Invar alloys[6,7,9,10]. However, such heat treatments have been predicted tolead to strength degradation owing to grain coarsening [15,16].Furthermore, embrittlement may occur after such a heat treatmentvia S grain-boundary segregation originating from saccharin[15,17–21].

The mechanical properties of the electrodeposited Invar alloyshave not been sufficiently explored, although the literaturecontains several reports [7–10] that address alloy compositionsand phases of electrodeposited Invar alloys. We previouslyreported the mechanical properties of electrodeposited Invar-likeFe-42mass%Ni containing a small amount of S [15]. However, themechanical properties of electrodeposited Fe–Ni alloys withcompositions in the Invar composition range, i.e., approximately36 mass% Ni, have not been studied with respect to the decrease intheir CTE upon heat treatment.

In this study, we prepared electrodeposited Invar alloys with Nicontents of 36, 38, and 40 mass% and evaluated their thermalexpansion and tensile properties. Moreover, we investigated theeffects of heat treatment on the alloys’ CTEs, microstructures, andmechanical properties, especially their ductility.

2. Experimental

Watts-type Ni plating bath (sulfate/chloride bath) was modifiedto prepare Fe-Ni alloy plating bath because a chloride bath exhibitshigher stress and a sulfamate bath exhibits risks of self-decomposition in an acidic media [22]. The plating bath comprisedFeSO4�7H2O, NiSO4�6H2O, and NiCl2�6H2O as metal sources, H3BO3

(boric acid), C7H4NNaO3S�2H2O (saccharin sodium dihydrate) as astress reducer, malonic acid. Malonic acid was used as a mask agentfor Fe3+ generated via electrochemical and air oxidation. Electro-deposited Fe–Ni alloys deposited in the absence of Malonic acidwill be brittle because of incorporation of Fe(OH)3 into the deposits[23]. Details of the bath composition and plating conditions arepresented in Table 1 [10]. To prepare the plating bath, we dissolvedeach reagent in distilled deionized water and performed anactivated-carbon treatment. The pH of the plating bath wasadjusted to 2.3 with 5% H2SO4 solution and NiCO3�2Ni(OH)2�4H2O.When the pH increases to 3 or higher, Fe3+ which cannot be

completely masked tends to precipitate as Fe(OH)3 althoughcurrent efficiency is higher. In contrast, in the case of pH of 2 orlower, current efficiency drastically decreases via substantialhydrogen generation [23]. In addition, malonic acid can act as apH buffering agent for pH 2 to 3 solutions because the pKa ofmalonic acid for the first ionization is 2.83 [24]. Galvanostaticelectrolysis was conducted at a current density of 40 mA/cm2 usinga constant-current power supply (YPP15100, YAMAMOTO-MS Co.,Ltd.). Lower current density (<10 mA/cm2) and higher currentdensity (approximately 100 mA/cm2) decrease current efficienciesvia proton reduction and water electrolysis, respectively. There-fore, current densities of approximately 40 mA/cm2 are suitable forhigher deposition rate, i.e., approximately 30 mm/h in this study. Astainless steel plate (SUS304) with a square exposed area of60 mm � 60 mm was used as a cathode, and a pure Fe sheet(YAMAMOTO-MS Co., Ltd.) and an electrolytic Ni sheet (SumitomoMetal Mining Co., Ltd.) inserted in anode bags were used as anodes.The plating bath was maintained at 50 �C in a 30-L electrolysis cell(YAMAMOTO-MS Co., Ltd.), and the electrolyte was filtrated andcirculated by a magnetic drive pump. Freestanding electro-deposited Invar Fe–Ni alloys were obtained as specimens bymechanical removal of the deposits from the stainless steelsubstrate. The thickness of the specimens was approximately100 mm. The composition of the electrodeposited Fe–Ni alloy wasdetermined by the fundamental parameter method using afluorescence X-ray analyzer (SEA6000VX, Hitachi High-TechScience Corp.). C and S content of the alloy were evaluated byan infrared absorption method using a carbon/sulfur analyzer(EMIA-V2, Horiba Ltd.).

The thermal expansion of the Invar alloys was measured by thecompressive load method using a thermal dilatometer (TD-5000S,NETZSCH Japan K. K.) with a load of 0.029 N. Specimens with alength of 20 mm and a cross-sectional area of 5 mm � 0.1 mm(width � thickness) were used for the thermal expansion measure-ments. The heating rate during the thermal expansion measure-ments was 5 �C/min, and the samples were maintained under aflowing nitrogen atmosphere (200 mL/min). The coefficient ofthermal expansion (CTE) was determined from the measurementsbetween 30 and 100 �C by the equation

CTE (�C�1) = 1/L (dL/dT), (1)

where L is the original length of the specimen.X-ray diffraction measurements of the Invar alloys were

performed using an X-ray diffractometer (Ultima IV, RigakuCorporation) equipped with a Cu Ka radiation source operatedat 40 kV and 40 mA.

Dumbbell-type specimens used for the tensile tests werefabricated by a photo-etching process using a ferric chloridesolution (35 �C). The dumbbell-type specimens were fabricatedwith a gauge section of 5 mm � 0.1 mm (width � thickness), agauge length of 10 mm, and an overall length of 50 mm, as shownin Fig. 1. The tensile tests were conducted using an Instron-typeuniversal tester (Autograph AG-250 kN, Shimadzu Corp.) under atensile speed of 0.5 mm/min and at room temperature. The

Fig. 1. Dimensions of the dumbbell-type specimens used for the tensile tests.

― 30 ᵒC ⇆ 200 ᵒC― 30 oC ⇆ 300 oC― 30 oC ⇆ 400 oC― 30 oC ⇆ 500 oC― 30 oC ⇆ 600 oC― Pyr ometall urgi call y

― 30 ᵒC ⇆ 200 ᵒC― 30 oC ⇆ 300 oC― 30 oC ⇆ 400 oC― 30 oC ⇆ 500 oC― 30 oC ⇆ 600 oC― Pyr ometall urgi call y

(a) ― 30 ᵒC ⇆ 200 ᵒC― 30 oC ⇆ 300 oC― 30 oC ⇆ 400 oC― 30 oC ⇆ 500 oC― 30 ᵒC ⇆ 600 ᵒC

(b) (c)

180 T. Nagayama et al. / Electrochimica Acta 205 (2016) 178–187

ultimate tensile strength (UTS) and elongation (d) were evaluatedas tensile properties of the alloys. To evaluate the UTSs and ds ofthe Invar alloys, maximum tensile loads were divided by the initialcross-sectional areas of the alloys and changes in gauge lengthswere divided by the initial gauge lengths of the specimens,respectively. The hardness of the Invar alloys was determinedusing a micro-Vickers hardness meter (HM-200, Mitutoyo Corp.)under an applied load of 0.49 N and at room temperature for centersections of the alloys with a thickness of approximately 100 mmmounted onto epoxy resins and mirror-polished. The specimenswere subjected to heat treatments at 300, 400, 500, and 600 �C for1 h under a vacuum of approximately 5 mPa using a vacuum-heat-treatment furnace (KDF-V50RM, Denken Co., Ltd.).

The microstructure and the fractured surface after the tensiletest of the Invar alloys were examined via field-emission scanningelectron microscopy (FE-SEM) on electron microscopes (JSM-6700F and JSM-7001F, JEOL Ltd.) equipped with an energy-dispersive X-ray spectroscopy (EDS) analyzer (JED-2300F, JEOL).The samples for the microstructure observations were mountedonto the epoxy resins and mirror-polished.

3. Results and Discussion

3.1. Composition, structure, and thermal expansion properties of theInvar Fe–Ni alloy electrodeposits

Fig. 2 shows the influence of FeSO4 added to the plating bathconsisting with the composition listed in Table 1 on the Fe contentof the electrodeposited Fe–Ni alloys. The Fe content increased withincreasing FeSO4 concentration; Fe-rich Fe–Ni alloys wereobtained at FeSO4 concentrations of 0.25 mol/L or greater. InvarFe–Ni alloys with Fe contents of 60 mass% (40 mass%Ni), 62 mass%(38 mass%Ni), and 64 mass% (36 mass%Ni) were prepared at FeSO4

concentrations of 0.3, 0.325, and 0.35 mol/L, respectively, consis-tent with our previously published results [10].

Saccharin has been widely used in the electrodeposition of Niand Fe–Ni alloys as an additive to reduce residual stress andimprove the roughness of electrodeposits [6–12,17,18]. However,

0

20

40

60

80

100

0 0.1 0.2 0.3 0.4

FeSO4 concentra�on /mol L-1

Fe c

onte

nt/m

ass%

Fig. 2. Influence of the addition of FeSO4 to the plating bath on the Fe content of theelectrodeposited Fe–Ni alloys.

saccharin promotes the codeposition of S into electrodeposited Niand Fe–Ni alloys [15,17,18]. Furthermore, it was reported that the Sis present as a metal sulfide (FeS, NiS), originating through theelectrochemical reduction of saccharin [25]. Our results confirmthat the electrodeposited Invar Fe–Ni alloys with 36–40 mass% Nicontain approximately 0.02 mass% S.

It is known that an organic additive addition often effect onimpurities, i.e. C, incorporation into Ni, Fe and their alloyselectrodeposits [26–34]. Amounts of incorporated impurities inthe deposits depend on types of organic additives. In addition,impurities in the electrodeposits closely related to mechanicalproperties such as hardness [26]. It was reported that effects oforganic acids on C contents and hardness of Fe electrodeposits [26].Although C contents and hardness of Fe electrodeposits depend ontypes of organic acid, effects of addition of malonic acid on Ccontents and hardness of Fe deposit were less [26]. In this study,there was no difference in C content in the electrodeposited Invaralloys between with and without malonic acid, �0.01mass%.Consequently, the C (�0.01mass%) in the Invar alloys wasapparently originated from saccharin [31–34].

An S (sulfide) in the electrodeposited Ni or Fe–Ni alloy easilyagglutinate together at grain boundaries, promoting grain-boundary embrittlement and intergranular fracture even duringannealing at low temperatures (�300 �C) [15,17–21]. Therefore,embrittlement of the electrodeposited Invar 36–40 mass%Ni alloyscontaining S might occur via grain-boundary segregation of S afterthe heat treatment. The effects of S codeposition on the mechanicalproperties of the alloys will be discussed later.

In our previous studies [6,10], the thermal expansion propertiesof Invar Fe-36 mass%Ni alloy and an Invar Fe–Ni alloy electro-formed metal mask were evaluated. These as-deposited Invar alloysamples exhibited negative CTEs in the temperature range from300 to 400 �C, suggesting thermal contraction of the alloys. In thisstudy, to investigate thermal expansion/contraction behaviors indetail, thermal expansion measurements were performed for theInvar Fe–Ni alloys with 36–40 mass% Ni.

Fig. 3 shows thermal expansion curves for the electrodepositedFe-36 mass%Ni alloy (a), Fe-38 mass%Ni alloy (b), and Fe-40 mass%Ni alloy (c) as a function of temperature, along with the curves forpyrometallurgically produced Fe-36 and 40 mass%Ni alloys. Inaddition, CTEs (from 30 to 100 �C) of the alloys after the precedingheat treatment at 300 to 600 �C during the thermal expansionanalysis are shown in Fig. 4 and Table 2. The CTEs of the as-deposited Fe-36 mass%Ni, Fe-38 mass%Ni, and Fe-40 mass%Nialloys were approximately 11, 10, and 9 ppm/�C, respectively,

0 100 200 300 400 500 6000 100 200 300 400 500 60 00 100 200 300 400 500 600

produ ced all oy

Ther

mal

exp

ansio

n, d

L/L

/%

0.2%

produ ced all oy

T /ᵒC

0.2%

T /ᵒC

0.2%

T /ᵒC

1st run 1st run 1st run

Pyr ometall urgi call yprodu cedFe-36 mass%N i all oy

Pyr ometall urgi call yprodu cedFe-40 mass%N i all oy

Fig. 3. Successive thermal expansion curves for the electrodeposited Fe-36 mass%Ni alloy (a), Fe-38 mass%Ni alloy (b), and Fe-40 mass%Ni alloy (c), along with resultsfor pyrometallurgically produced Fe-36 mass%Ni and Fe-40 mass%Ni alloys.

0

2

4

6

8

10

12

0 100 200 300 400 500 600

CTE

(30–

100°

C) /

ppm

ᵒC-1

Heat-treatment temperature /ᵒC

Fig. 4. Effects of heat-treatment temperature on the CTEs of the electrodepositedInvar Fe–Ni alloys with Ni contents of 36, 38, and 40 mass%. ~ 36 mass%Ni, �38 mass%Ni, * 40 mass%Ni. Fig. 5. X-ray diffraction patterns of the electrodeposited Invar Fe–Ni alloys with Ni

contents of 36, 38, and 40 mass% with and without heat treatment at 300 to 600 �C.! fcc, 5 bcc.

T. Nagayama et al. / Electrochimica Acta 205 (2016) 178–187 181

which are 2 to 10 times larger than those of the pyrometallurgicallyproduced Invar Fe–Ni alloys (1 to 4 ppm/�C [1,2]). The CTEs of theelectrodeposited Fe–Ni alloys increased with decreasing Nicontent. This result is contrary to those obtained for pyrometallur-gically produced Invar Fe–Ni alloys with Ni contents of 36, 38, and40 mass% (CTEs of 1, 2, and 4 ppm/�C, respectively). For all of theFe–Ni alloy samples, thermal contractions originated fromstructural change were observed at 300 to 400 �C. In the case ofthe Fe-40 mass%Ni alloy, the temperature at which contractionbegan was lower, i.e., 200 �C, than those for the Fe–Ni alloys withlower Ni contents. In contrast, for pyrometallurgically producedInvar Fe–Ni alloys with 36–40 mass% Ni, there was no thermalcontraction. Those pyrometallurgically produced Invar Fe–Nialloys exhibited continuous thermal expansion. In the case ofthe alloys heat-treated at 400 �C, the CTEs drastically decreased toapproximately 5 ppm/�C. Furthermore, the CTEs reached valuessimilar to those of pyrometallurgically produced Invar Fe–Ni alloysheat-treated at 500 �C or higher. In our preliminary work [10], wereported that the thermal contraction and the CTE variation of theFe-36 mass%Ni alloy with heat treatment correspond with changesin the microstructure of the alloy. The phenomenon of thedimensional variation with the preceding heat treatment in thecase of the Fe–Ni alloys with 36–40 mass% Ni will be discussed indetail later. The extent of contraction after the heat treatment to600 �C increased with increasing Fe content of the Invar Fe–Nialloys. The moduli of the contraction after heat treatment to 600 �Cfor Fe-36 mass%Ni, Fe-38 mass%Ni, and Fe-40 mass%Ni alloys were0.34, 0.33, and 0.31%, respectively.

Fig. 5 shows X-ray diffraction (XRD) patterns of the electro-deposited Invar Fe-36 to 40 mass%Ni alloys with and without heattreatment at 300 to 600 �C. The as-deposited Invar Fe–Ni alloyswith 36–40 mass% Ni were mainly composed of bcc phases with(100) orientation. This composition differs from the equilibriumfcc phase [13] produced pyrometallurgically, as already reported[7–10]. Only in the case of Fe-40 mass%Ni alloy, a weak fccdiffraction peak was detected in the as-deposited state. Thus far,the metastable bcc phases favored in the electrodeposited Fe–Nialloy with Invar composition range were observed in numerous

Table 2Coefficients of thermal expansion (CTEs) of the electrodeposited Invar Fe–Ni alloys wit

Heat-treatment temperature/�C CTE (30 to 100 �

Fe-36 mass%Ni

As-deposited 10.6

300 9.7

400 4.1

500 3.3

600 1.7

Pyrometallurgically produced Invar Fe–Ni alloys [2] 1.2

investigations [7–10,35]. Vicenzo [35] suggested that hydrogencodeposition through the intermediate hydrides formation[35–37] might have a role in promoting the bcc phase formationin the electrodeposited Fe–Ni alloy with Invar composition range,although the details are not disclosed. In our electrodeposition ofInvar alloys, because a current efficiency was approximately 70%due to hydrogen revolution reaction, we supposed that thehydrogen revolution reaction might play a role as a factorinfluencing the metastable bcc phase formation of our Invaralloys. The phase of the electrodeposited Fe-36 mass%Ni alloys isknown to depend on the experimental conditions. Our as-deposited Invar alloys had (100)-oriented bcc phases, whereasthe bcc phases exhibited (110) preferred orientation in the case ofthe electrodeposited Invar alloys reported by Grimmett et al. [7]and Tabakovic et al. [9], which were obtained at lower currentdensity (5–26 mA/cm2) than that of our Invar electrodeposition. Inaddition, for as-deposited Invar alloys reported by Kim et al. [8],(110)-oriented bcc phases were obtained from baths with lowersaccharin (1.5–3 g/L). In contrast, (100)-oriented bcc phases wereobtained from baths with higher saccharin (4.5–6.5 g/L). In the caseof our alloys, after the 300 �C heat treatment, the intensities of thebcc peaks decreased, whereas those of the fcc peaks increased.These results indicate that a transformation from metastable bccphases to equilibrium fcc phases was initiated at 300 �C.Furthermore, after heat treatment at 500 �C or higher, the XRDpatterns of the electrodeposited Fe–Ni alloys with Ni contents of36, 38, and 40 mass% indicated a single fcc phase with latticeconstants of 0.3592, 0.3595, and 0.3594 nm, respectively. Thelattice constants of the alloys are similar to those of pyrometallur-gically produced Fe–Ni alloys. The lattice constants of pyrome-tallurgically produced Fe–Ni alloys with Ni contents of 35.1, 38.4,and 40.7 mass% are 0.3593, 0.3596, and 0.3596 nm [38],respectively. The diffraction peaks in the XRD pattern of the InvarFe–Ni alloys prepared by annealing at 600 �C were sharp.

In the case of the Invar Fe-36 mass%Ni alloys composed ofalmost only bcc phase in the as-deposited state, the calculated

h Ni contents of 36, 38, and 40 mass%.

C)/ppm �C�1

alloy Fe-38 mass%Ni alloy Fe-40 mass%Ni alloy

9.8 8.69.2 8.54.6 5.44.0 5.12.8 4.42.2 3.9

0200400600800

10001200

0 100 200 300 400 500 600

0

10

20

30

40

0 100 200 300 400 500 600

100

200

300

400

0 10 0 200 30 0 40 0 500 600

Ul�

mat

e te

nsile

stre

ngth

,

UTS

/MPa

Vic

kers

har

dnes

s, H

VEl

onga

�on,

δ/%

Heat treatment temp erature /ᵒC

Fig. 7. Effects of temperature of heat treatment on the ultimate tensile strength(UTS), elongation (d), and hardness of the electrodeposited Invar Fe–Ni alloys withNi contents of 36, 38, and 40 mass%. ~ 36 mass%Ni, � 38 mass%Ni, * 40 mass%Ni.

Fig. 6. Stress–displacement curves for the electrodeposited Invar Fe–Ni alloys withNi content of 36, 38, and 40 mass% with and without heat treatment at 300 to600 �C.

182 T. Nagayama et al. / Electrochimica Acta 205 (2016) 178–187

modulus of the contraction after the sample was annealed at600 �C was 0.35%, assuming transformation from the bcc phase tothe fcc phase with lattice constants of 0.2861 and 0.3592 nm,respectively. This modulus is comparable with the value of 0.34%obtained from the thermal expansion measurements reported inFig. 3. This similarity demonstrates that the thermal contraction ofthe Invar Fe-36 mass%Ni alloy is attributable to an increase of theatomic packing density by the phase transformation from a bccphase to an fcc phase, as pointed out in our preliminary work [10].We considered that the thermal contraction phenomena of theInvar Fe–Ni alloys with 38 and 40 mass% Ni were also generated bya similar mechanism. These results show that the as-depositedInvar Fe-36 to 40 mass%Ni alloys with CTEs 2 to 10 times largerthan those of pyrometallurgically produced Fe–Ni alloys weremainly composed of bcc phases. In the samples annealed at 400 �Cor higher, the fcc phases became the main phases, accompanied bya drastic decrease of the CTEs of the alloys (Fig. 4), consistent withpreviously reported results [7,9,10]. Tabakovic et al. [9] alsoindicated that a heat treatment at 400 �C or above was necessary toobtain normal thermal expansion properties corresponding to themagnetic properties of Invar Fe-36 mass%Ni alloy. In addition,ferromagnetic bcc Invar Fe-36 mass%Ni alloy with a high saturationmagnetization (Bs) has been reported to transform to antiferro-magnetic fcc Invar alloy with a low Bs upon treatment at 400 �C orabove [9]. Moreover, the low thermal expansion properties ofpyrometallurgically produced Invar Fe–Ni alloys with 30 to40 mass% Ni are known to originate from the alloys’ antiferromag-netic property, i.e., a low Bs [2,39]. Consequently, the change in themain phase of the electrodeposited Invar Fe-36 to 40 mass%Nialloys from bcc to fcc by annealing at 400 �C suggests that the CTEsof the alloys drastically decreased because of the magnetic-property transition from ferromagnetic to antiferromagnetic, asreported by [9]. We considered that CTEs of the electrodepositedInvar Fe–Ni alloys with 36 to 40 mass% Ni corresponded to the fcc/bcc volume fraction in the alloys. In particular, the CTEs of thealloys approached the values of the produced pyrometallurgicallyalloys with increasing fraction of the equilibrium fcc in the alloys,as previously reported [10].

3.2. Mechanical properties of the invar Fe–Ni alloy electrodeposits

Fig. 6 shows representative stress-displacement curves of theelectrodeposited Invar Fe-36 to 40 mass%Ni alloys with andwithout heat treatment at 300 to 600 �C. Moreover, the effects ofthe heat-treatment temperature on the UTS and d determinedfrom these tensile tests and on the hardness of the electro-deposited Invar Fe–Ni alloys with 36 to 40 mass% Ni are shown inFig. 7 and Table 3, respectively. The UTS and hardness of the as-deposited Invar Fe–Ni alloys with 36 to 40 mass% Ni exhibitedapproximately 700 MPa and 250HV, respectively. These UTS andhardness values of the alloys are lower than those of theelectrodeposited Fe-80 mass%Ni alloy (1760 MPa, 550HV) [15],Fe-42 mass%Ni alloy (1610 MPa, 430HV) [15], and Ni (1010 MPa,340HV) [15] prepared in plating baths containing saccharin as wellas these Invar alloys. However, the UTS and hardness values of theas-deposited Invar alloys were higher than those of conventionalInvar alloys, whose UTS and hardness values are approximately500 MPa and 150HV, respectively [40]. The d of the as-depositedInvar Fe–Ni alloys was approximately 10%, similar to that ofconventional electroformed Ni for an electroforming application[19,41]. This result indicates that the as-deposited Invar alloysexhibit good ductility. The UTS and hardness slightly increased,and the d decreased to some extent after heat treatment at300–400 �C. The high values of UTS and hardness and the goodductility of the Invar alloys remained after heat treatment attemperatures up to 500 �C. When the annealing temperature was

Table 3Effects of heat treatment temperature on ultimate tensile strength (UTS), elongation (d), and hardness of Invar Fe–Ni alloys with Ni contents of 36, 38, and 40 mass%.

Heat treatment temperature/�C UTS/MPa

Fe-36 mass%Ni alloy Fe-38 mass%Ni alloy Fe-40 mass%Ni alloy

As-deposited 692 704 730300 770 823 915400 812 854 871500 719 720 805600 479 551 560Pyrometallurgically produced Invar Fe-36mass%Ni alloy [40] 517 (annealed)

Heat treatment temperature/�C d/%

Fe-36 mass%Ni alloy Fe-38 mass%Ni alloy Fe-40 mass%Ni alloy

As-deposited 12.9 15.3 9.1300 9.6 6.9 4.8400 7.7 7.0 2.6500 12.5 9.9 5.8600 29.5 27.7 23.4Pyrometallurgically produced Invar Fe-36mass%Ni alloy [40] 30 (annealed)

Heat treatment temperature/�C Hardness/HV

Fe-36 mass%Ni alloy Fe-38 mass%Ni alloy Fe-40 mass%Ni alloy

As-deposited 243 241 274300 260 272 328400 261 269 281500 222 246 269600 139 190 173Pyrometallurgically produced Invar Fe-36mass%Ni alloy [40] 120 (annealed)

T. Nagayama et al. / Electrochimica Acta 205 (2016) 178–187 183

increased to 600 �C, the UTS and the hardness sharply decreased,accompanied by an increasing of d. The UTS, hardness, and d of theInvar alloys have values similar to those of conventional as-castInvar alloys (�400 MPa, 120HV, and 30%, respectively [40]). Adrastic decrease in the d value to less than 1% has been reported tooccur at 400 �C or above for electrodeposited Ni [15,17] and Fe–Nialloys with Ni contents of 60 mass% or above [15,18,21] andcontaining a small amount of S. Notably, in the present work, the Scontent of the electrodeposited Invar alloys was approximately0.02 mass%; nevertheless, no severe embrittlement was observedafter heat treatment at 300 to 600 �C. The Invar Fe–Ni alloys withhigher Ni contents tended to exhibit higher strength. In addition,the UTS was almost directly proportional to the hardness, whereasd was almost inversely proportional to strength.

Fig. 8 shows FE-SEM images of fractured faces of tensile testspecimens of the electrodeposited Invar Fe-36 mass%Ni alloy withand without heat treatment at 400 and 600 �C. Dimple patternsindicating plastic deformation accompanied by good ductility wereobserved for all Fe–Ni alloy specimens. The fractured face of the as-

Fig. 8. FE-SEM images of the fractured faces of representative tensile test specimens of t400 and 600 �C. The insets show photos of the specimens after tensile fracture.

deposited alloy consisted of dimples with a diameter ofapproximately 1 mm. Finer dimples of approximately 0.5 to1 mm were also observed after heat treatment at 400 �C, whereelongation of the alloy slightly decreased, corresponding to slightstrengthening. Moreover, markedly large dimples with a diameterof approximately 3 mm were observed after the heat treatment at600 �C, where the d substantially increased, as shown in Figs. 6 and7. In addition, after heat treatment at 400 to 500 �C, the strengthremained relatively high, whereas the CTE decreased drastically;the heat treatment resulted in high strength, compatible with thelow CTE of the alloys. Tensile-fractured faces characterized byintergranular cleavage have been reported for S-containing Ni andNi-rich Ni–Fe alloys after heat treatment, where the samplesexhibited brittle behavior with extremely low d (<1%) values[15,17,18]. We confirmed that the electrodeposited Invar Fe–Nialloys exhibited ductile behavior, irrespective of whether thesamples were heat-treated.

Fig. 9 shows backscattered-electron images of the cross sectionof the electrodeposited Invar Fe- 36 to 40 mass%Ni alloys with and

he electrodeposited Invar Fe-36 mass%Ni alloy with and without heat treatment at

Fig. 10. Backscattered-electron image (BEI) and EDX mapping images of the electrodeposited Invar Fe-36 mass%Ni alloy heat-treated at 600 �C.

Fig. 9. Backscattered-electron images of microstructures of the electrodeposited Invar Fe–Ni alloys with and without heat treatment at 300 to 600 �C. Black arrows show thefilm growth direction.

184 T. Nagayama et al. / Electrochimica Acta 205 (2016) 178–187

Fig. 11. Backscattered-electron image (BEI) and EDX spectra of a matrix point (a)and a precipitate point (b) on the electrodeposited Invar Fe-36 mass%Ni alloy heat-treated at 600 �C. The cross marks in the BEI indicate the EDX measurement points.

T. Nagayama et al. / Electrochimica Acta 205 (2016) 178–187 185

without heat treatment at 300 to 600 �C. The electrodepositedInvar alloys without heat treatments exhibited a columnarmicrostructure composed of micrometer-sized coarse grains andsubmicrometer-sized grains without defects such as cracks andvoids. In addition, the electrodeposited Invar alloys were mainlycomposed of the bcc phase with (100) preferred orientation, asalready discussed (Fig. 5). These microstructure characteristics arecorrelated to the addition of saccharin to the plating baths, as notedby Kim et al. [8].

The microstructure of the alloys with higher Ni contents tendedto comprise a larger fraction of the submicrometer-sized finegrains. In addition, the columnar microstructures became increas-ingly evident with decreasing Ni content of the alloys. After theheat treatment at 300 to 400 �C, where the proportion of fcc phaseincreased (Fig. 5), the columnar structure became indistinct andthe fraction of the submicrometer-sized fine grains increased. Heattreatment at 500 �C led to grain growth; the alloys treated at thistemperature were fully composed of the fcc phase, accompaniedby a decrease in the CTE, as shown in Figs. 4 and 5. However,extensive grain growth was not confirmed for the sample heat-treated at 500 �C. This result suggests that the higher UTS andhardness values of the alloys (Figs. 6 and 7) remained unchangedupon heat treatment 500 �C because of these fine microstructures.The significant grain growth to several micrometers in size causeda decrease in the UTS and hardness and an increase in the d of theelectrodeposited Invar alloys upon heat treatment at 600 �C(Figs. 6 and 7) [42]. In addition, precipitates were observed inthe alloys heated at 600 �C, as indicated by arrows in Fig. 9.

Fig. 10 shows a backscattered-electron image and EDX mappingimages of the Invar Fe-36 mass%Ni alloy after heat treatment at600 �C. Granular precipitates with various sizes ranging fromapproximately 0.1 to 0.5 mm formed in grain-boundary triple-points or were entrapped in crystal grains of the alloy. Fig.11 showsthe results of the EDX chemical analysis at matrix grain (a) and atprecipitate (b) in the Invar Fe-36 mass%Ni alloy heat-treated at600 �C. No substantial S Ka X-ray signal was detected at the Fe–Nialloy matrix grain, as shown in the EDX spectra of point (a). Thechemical composition of the Fe–Ni alloy matrix was 59 mass%Fe–41 mass%Ni (60at%Fe–40at%Ni), as calculated from the EDX resultswith the carbon signals generated by contamination excluded. Thiscalculated value is comparable with that of the bulk composition ofthe alloy, 64 mass%Fe–36 mass%Ni, measured by fluorescence X-ray analysis. By contrast, at precipitate point (b), intense Fe and Ssignals were detected. Because a Ni X-ray spectrum was notobserved, the Ni content of the precipitate was apparently verylow. The chemical composition of the precipitate point was75 mass%Fe–25 mass%S (64at%Fe–36at%S), as calculated on thebasis of the intensity of Fe and S X-ray peaks. This S content issubstantially higher than the value of 6.3 mass% S reported byWang et al. [17] to have segregated to the Ni grain boundaries. Ourresults suggest that S in the Invar Fe-36 mass%Ni alloy becamehighly concentration upon heat treatment at 600 �C and formedgranular Fe–S compounds or Fe-rich Fe–Ni–S in the matrix grainsor grain-boundary triple-points. The precipitates observed in Fe-38 mass%Ni and Fe-40mass%Ni alloys heat-treated at 600 �C,shown in Fig. 9, are similar to those observed in Fe-36mass%Nialloys.

Electrodeposited Ni and Ni-rich Fe–Ni alloys containing smallamounts of S are often embrittled upon annealing. The d of theelectrodeposited Ni containing 0.046 mass% S was drasticallydecreased from 3% to less than 0.3% in samples annealed at 300 �Cor higher temperatures [17]. The electrodeposited Fe-80 mass%Nialloy with 0.2 mass% S was embrittled by annealing at 400 to700 �C: the d decreased from 7% to less than 1% [18]. An S (sulfide)in the as-deposited fcc films is dispersed at grain boundaries asreported by Tabakovic [25]. During recrystallization and grain

growth of fcc grains by heat treatment, the S (sulfide) is extrudedinto recrystallized fcc grain boundary with segregating as a thinlayer form. The S thin layer coats or wets the grain boundaries andcauses grain boundary weakening; therefore, distinct embrittle-ment occurs [18]. Furthermore, the embrittlement is promoted byrapid grain growth during annealing [17,18]. In contrast, an S(sulfide) in our as-deposited bcc Invar alloys will also dispersed atbcc grain boundaries. During grain growth of transformed fcc grain(equilibrium phase) by heat treatment, the S (sulfide) segregatedas a granular sulfide form at primary bcc grain boundaries. Upon

186 T. Nagayama et al. / Electrochimica Acta 205 (2016) 178–187

heat treatment at 500 �C or above, bcc grain eliminated accompa-nied by fcc grain growth. Consequently, the granular sulfideagglutinated furthermore and was entrapped in the matrix grainsor grain-boundary triple points of transformed fcc grains. Theelectrodeposited Invar Fe–Ni alloys did not exhibit any brittleintergranular cleavage but did show ductile fracture, as shown inFig. 8, even though it contained 0.02 mass% S. Furthermore, the d ofthe alloys was slightly decreased by annealing at temperatures upto 400 �C but remained larger than 3%. These values are muchlarger than those reported for annealed samples of Ni [17] andFe-80 mass%Ni alloy [18] containing S, which were less than 1%.After heat treatment at 500 �C or above, the grains grew and the dincreased to approximately 30% at 600 �C. The S in the Invar Fe–Nialloys heat-treated at 600 �C existed not as a filmy sulfide at grainboundaries, which causes grain boundary embrittlement of Ni andNi-rich Fe–Ni alloys but as a granular sulfide that prevented Ssegregation to grain boundaries, thereby preventing grain-boundary embrittlement. Moreover, it is considered that thesetwo-type morphologies of the sulfide, i.e. as a coating at grainboundaries or as a precipitate, will be determined by grain growthform during an annealing, i.e. from primary fcc grain to recrystal-lized fcc grain, or from primary bcc to transformed fcc in ourresults, respectively.

4. Conclusions

1. Electrodeposited Invar Fe–Ni alloys with 36 to 40 mass% Ni wereprepared from plating baths containing saccharin as a stressreducer and containing various Fe2+ concentrations. The InvarFe–Ni alloys contained of small amount of S (�0.02 mass%).

2. The CTEs of the as-deposited Invar Fe–Ni alloys wereapproximately 9 to 11 ppm/�C and were larger than those ofpyrometallurgically produced Invar alloys. When the alloyswere heat-treated at 400 to 500 �C, their CTEs drasticallydecreased to approximately 5 ppm/�C. Furthermore, upon heattreatment at 600 �C, the CTEs reached approximately 2 to4 ppm/�C depending on alloy composition; these CTEs arecomparable with those of pyrometallurgically produced alloys.

3. The as-deposited Invar Fe–Ni alloys were mainly composed ofmetastable bcc phases, resulting in larger CTEs. When the alloyswere annealed at 400 �C or above, the equilibrium fcc phasesbecame the predominant phases, accompanied by a drasticdecrease of the CTEs. The bcc-to-fcc transformation led to adecrease of the CTEs and to thermal contractions.

4. Upon the heat treatment, an S (sulfide) at bcc grain boundariessegregated not as a thin film but as a granular sulfide form atprimary bcc grain boundaries in the electrodeposited Invaralloys. Upon heat treatment at 500 �C or above, bcc graineliminated accompanied by fcc grain growth and the granularsulfide agglutinated further. Consequently, the agglutinatinggranular sulfide was entrapped in the matrix grains or grain-boundary triple-points of transformed fcc grains. In addition, itis considered that these two-type morphologies of the sulfide,as a thin film at grain boundaries or as a precipitate, will bedetermined by grain growth form during an annealing.

5. Ductile behavior of the electrodeposited Invar Fe–Ni alloys wasconfirmed, irrespective of whether the alloys were heat-treated.Upon heat treatment at 400 to 500 �C, the Invar Fe–Ni alloysexhibited high strength with good ductility, consistent withtheir low CTE. After the heat treatment, no severe embrittle-ment of the electrodeposited Invar alloys was observed despitethe codeposition of S because the S existed as a granular sulfide,thereby preventing grain-boundary embrittlement.

Acknowledgement

This work was supported by the Program for Fostering RegionalInnovation Kyoto Environmental Nanotechnology Cluster of TheMinistry of Education, Culture, Sports, Science and Technology,Japan.

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