Acta Biomaterialia 7 (2011) 1387–1397

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Acta Biomaterialia

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Hydrogen release from titanium hydride in foaming of orthopedic NiTi scaffolds

Shuilin Wu a,b, Xiangmei Liu a,b,c, K.W.K. Yeung c, Tao Hu b, Zushun Xu a,b, Jonathan C.Y. Chung b,Paul K. Chu b,⇑a Faculty of Materials Science and Engineering, Hubei University, Wuhan 430062, People’s Republic of Chinab Department of Physics and Materials Science, City University of Hong Kong, Tat Chee Avenue, Kowloon, Hong Kongc Division of Spine Surgery, Department of Orthopaedics and Traumatology, The University of Hong Kong, Pokfulam, Hong Kong

a r t i c l e i n f o

Article history:Received 26 June 2010Received in revised form 6 September 2010Accepted 12 October 2010Available online 20 October 2010

Keywords:Hydrogen releaseFoamsPorous NiTiCreep expansionScaffold

1742-7061/$ - see front matter � 2010 Acta Materialdoi:10.1016/j.actbio.2010.10.008

⇑ Corresponding author. Tel.: +852 34427724; fax:E-mail address: paul.chu@cityu.edu.hk (P.K. Chu).

a b s t r a c t

Titanium hydride powders are utilized to enhance the foaming process in the formation of orthopedicNiTi scaffolds during capsule-free hot isostatic pressing. In order to study the formation mechanism,the thermal behavior of titanium hydride and hydrogen release during the heating process are system-atically investigated in air and argon and under vacuum by X-ray diffraction (XRD), thermal analysis,including thermogravimetric analysis and differential scanning calorimetry, energy dispersive X-rayspectroscopy, and transmission electron microscopy. Our experiments reveal that hydrogen is continu-ously released from titanium hydride as the temperature is gradually increased from 300 to 700 �C.Hydrogen is released in two transitions: TiH1.924 ? TiH1.5/TiH1.7 between 300 �C and 400 �C and TiH1.5/TiH1.7 ? a-Ti between 400 �C and 600 �C. In the lower temperature range between 300 �C and 550 �Cthe rate of hydrogen release is slow, but the decomposition rate increases sharply above 550 �C. TheXRD patterns obtained in air and under vacuum indicate that the surface oxide layer can deter hydrogenrelease. The pressure change is monitored in real time and the amount of hydrogen released is affected bythe processing temperature and holding time. Holding processes at 425 �C, 480 �C, 500 �C, 550 �C, and600 �C are found to significantly improve the porous structure in the NiTi scaffolds due to the stepwiserelease of hydrogen. NiTi scaffolds foamed by stepwise release of hydrogen are conducive to the attach-ment and proliferation of osteoblasts and the resulting pore size also favor in-growth of cells.

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1. Introduction

Since the first report on the use of metal plates and screws infixing bone fractures in 1890 [1] artificial implants have beenwidely adopted in the repair and replacement of injured bone. Incomparison with dense materials, porous materials have a numberof advantages, because their interconnected porous structure al-lows in-growth of new tissues and the passage of nutrients. Fur-thermore, the porous structure typically yields a reduced Young’smodulus closer to that of human bone while the scaffold can stillmaintain enough strength to satisfy load-bearing demands [2,3].In particular, three-dimensional (3-D) microporous NiTi scaffoldshave immense potential in surgical implants due to the porous nat-ure of the materials and the advantageous properties inheritedfrom dense NiTi shape memory alloys, such as good biocompatibil-ity, as well as the shape memory and super-elasticity effects [4–9].A hierarchical nanostructure which can be produced on the ex-posed surface of 3-D microporous Ti-based scaffolds by chemicalmeans to match the organization of natural bone at the lowest

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level favors new bone tissue formation and fixation of bone im-plants [10].

The manufacture and application of porous metals have recentlybeen reviewed by Banhart [11] and porous NiTi has been producedby various powder metallurgical (PM) methods [12–17]. In thisprocess high melting point metals and alloys are foamed byutilizing creep expansion of a prefilled gas at elevated temperature[18–21]. The foaming agent creates the space within the structureand so choice of the right foaming agent is crucial to the process.Various materials, such as NH4HCO3 [22,23], NaCl [24], carbamideparticles [23], magnesium [25], saccharose, and poly(methyl meth-acrylate) [26], have been employed as space holders/foamingagents in PM processes to fabricate 3-D microporous NiTi/Tiscaffolds. Titanium hydride is also often used in the foaming oflow melting point metals such as aluminum and aluminum alloysbecause a large amount of hydrogen is released during heating[27–31]. Although titanium hydride have recently been used tofoam high melting point metals such as Ti and NiTi [18,32–34],there has not been a systematic investigation into the mechanismand role of the thermal evolution of titanium hydride pertainingto NiTi which constitutes a complex system due to self-reactionbetween the two elements at elevated temperature. Our previous

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1388 S. Wu et al. / Acta Biomaterialia 7 (2011) 1387–1397

results reveal the existence of trace amounts of titanium oxide onthe exposed surface during the capsule-free hot isostatic pressing(CF-HIP) process [22] and oxygen even under vacuum conditionsand high pressure argon backfilling. A similar problem will beencountered when using titanium hydride as the foaming agent,and so it is crucial to understand the thermal evolution of titaniumhydride in air and under vacuum. We here present the relationshipbetween hydrogen release and temperature, as well as holdingtime, by real time monitoring of the pressure. Based on the knowl-edge acquired and a better understanding of the mechanism weare able to produce NiTi scaffolds using titanium hydride as thefoaming agent.

2. Experimental procedures

2.1. Powder treatment

Nickel and titanium powders less than 75 lm in size were pur-chased from Shanghai Reagent Corporation and TiH2 powders witha nominal size of 640 lm were supplied by Beijing Research Insti-tute for Nonferrous Metals. The powders had a purity of 99.5%. Tostudy the thermal evolution in an oxygen atmosphere titanium hy-dride powders were treated in air in muffle furnaces at tempera-tures of 300 �C, 350 �C, 400 �C, 450 �C, 500 �C, 550 �C, 600 �C,800 �C, and 1000 �C for 0.5 h. The thermal evolution behavior oftitanium hydride was also investigated under vacuum from350 �C to 1000 �C in a chemical vapor deposition (CVD) system(Lenton Thermal Designs, London) at an initial pressure of3.5 � 10�4 Torr. After setting the temperature, the vacuum cham-ber was evacuated using a diffusion pump for about 45 min andthen backfilled with argon (flow rate of 5 sccm) to minimize oxida-tion while both a mechanical and the diffusion pump continued towork to maintain the vacuum during the process. Titanium hydridepowders (2.36 g) were placed in a ceramic crucible. The heatingrate was typically 30 �C min�1 and the samples were treated for0.5 h at the selected temperature and cooled to room temperaturein the furnace.

2.2. NiTi foams preparation

The NiTi foams were prepared by CF-HIP under argon using twotypes of sintered powders: (1) nickel and titanium and (2) nickeland titanium hydride having the same atomic ratio of nickel totitanium. The former is termed ‘‘control foam”. These powderswere thoroughly mixed in a horizontal universal ball mill andpressed into green compacts at a pressure of 200 MPa using ahydraulic machine before sintering by CF-HIP. More details aboutthe mixing, cold pressing, and CF-HIP process can be found else-where [22,35,36]. The control NiTi foam was produced by the gen-eral CF-HIP process, but the process to produce NiTi foams fromnickel and titanium hydride powders were different. We adoptedtwo processes according to the hydrogen release from titanium hy-dride. In the first process (process 1) the samples were heated to425 �C, 480 �C, 500 �C, 550 �C, and 600 �C for 30, 30, 20, 20, and10 min, respectively. In the other process (process 2) the conven-tional process in which the system was heated to the designedtemperature directly was adopted.

2.3. Thermal analysis

To investigate the thermal evolution of titanium hydride as afunction of temperature thermal analyses were performed to mon-itor weight loss and temperature in TGA-Q50 and TA SDT Q600thermogravimetric analyzers. The TGA-Q50 was equipped with asilicon carbide furnace as well as a platinum pan connected to an

electrical balance. The platinum pan was filled with titanium hy-dride powder and inserted into the furnace, which was backfilledwith a continuous flow of helium at a rate of 100 ml min�1. Itwas heated from 30 �C to 1000 �C at a rate of 10 �C min�1. In orderto reduce the variability in the thermogravimetric analysis (TGA)and study the phase transition of titanium hydride during TGA,the same thermogravimetric analyzer was used to repeat theexperiments from 30 �C to 600 �C with a holding time of 2 h at600 �C. Here, the initial weight of titanium hydride was17.073 mg. After TGA the phase composition of the remainingpowders was determined by XRD. Owing to limitations of theTGA-Q50 the TG analysis was performed in argon in a TA SDTQ600 at a flow rate of 100 ml min�1.

The phase transition change in the titanium hydride with tem-perature was determined by differential scanning calorimetry(DSC) (Setsys 16/18, Setaram Scientific and Industrial Equipment).Titanium hydride powder (14.2 mg) was placed in an alumina cru-cible inside the furnace which was backfilled with argon prior toheating under a continuous argon flow (50 ml min�1). Afterwards,it was heated from 200 �C to 840 �C at a rate of 10 �C min�1.

2.4. X-ray diffraction

XRD was conducted on a Siemens D500 X-ray diffractomer witha CuKa X-ray source operated at 40 kV and 30 mA. The powdersheated at 800 �C and 1000 �C were ground into small particles be-fore the XRD measurements because the powders had been sin-tered into large particles or clusters at the higher temperature.The spectra were acquired in the 2h range 20�–80� at a step incre-ment of 0.05�. The XRD patterns were studied based on the JCPCstandard diffraction database using the software MDI JAD5.0.

2.5. Real time pressure monitoring

By means of high vacuum CVD (described in Section 2.1) the re-lease of hydrogen as a function of temperature and time was eval-uated by monitoring the changes in the pressure in real time usinga vacuometer.

2.6. Powders and NiTi foams analysis

The particle size and morphology of the powders were exam-ined by scanning electron microscopy (SEM) (JSM5200 andJSM820). The elemental contents of the powders were determinedby energy dispersive X-ray spectroscopy (EDS) (average of fivemeasurements from different areas). The microstructure and com-position were characterized by transmission electron microscopy(TEM) and selected area electron diffraction (SAED). The micro-structure of the optimized NiTi foam was observed by SEM. Thegeneral porosity of the samples was determined from

e ¼ 1� mqV

� �� 100ð%Þ, where m and V are the mass and volume

of the porous samples, respectively and q is the theoretical densityof NiTi (6.45 g cm�3 for bulk equiatomic NiTi). The open porosity isthe percentage of open pores to total pores determined accordingto the ASTM standard B328-96 protocol [37]. The NiTi foam platesused in the microstructure analysis were cut from the sinteredfoams using a linear cutting machine, mechanically polished withgrit paper progressively up to 800 mesh, rinsed with alcohol, andoven dried.

2.7. Cytotoxicity evaluation

The samples used for cell culture were linearly cut from the sin-tered NiTi bars into disks 2 mm thick and 5 mm in diameter. Thesedisks were mechanically polished, ultrasonically washed three

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S. Wu et al. / Acta Biomaterialia 7 (2011) 1387–1397 1389

times with acetone and deionized water, and then sterilized in anautoclave at 120 �C for 30 min. Osteoblasts isolated from calvarialbones of 2-day-old mice that ubiquitously expressed an enhancedgreen fluorescent protein (EGFP) were cultured in a Dulbecco’smodified Eagle’s medium (DMEM) (Invitrogen) supplemented with10 vol.% fetal bovine serum (Biowest, France), antibiotics (100 Uml�1 penicillin and 100 lg ml�1 of streptomycin), and 2 mM L-glu-tamine at 37 �C in an atmosphere of 5% CO2 and 95% air. The sam-ples were affixed to the bottom of a 24-well tissue culture plate(Falcon) using 1% (w/v) agarose. A cell suspension consisting of15,000 cells in 100 ll of medium was seeded onto the samples.Wells without any metal disks served as controls. The cells were al-lowed to settle and attach to the surface of the disks for 4 h. 1 ml ofmedium was added to the wells, which was changed every 3 days.Four samples of each type were tested to improve the statistics andcell proliferation was evaluated after culture for 8 days. The cellswere allowed to attain confluence during the examination period.The morphology of the attached living EGFP-expressing osteoblastswas examined by SEM after the disks with attached cells werewashed gently in PBS before fixing in 2.5% glutaraldehyde bufferedat a pH of 7.42 with 0.1 M sodium cacodylate for 24 h at 4 �C. Theexcess contents were removed by washing in 0.1 M sucrose in cac-odylate buffer and dehydration was carried out in a series of bathscontaining different ratios of ethanol to water up to 100% ethanol.The cells were then critical point dried and sputter coated withgold before examination by SEM (JSM-820).

3. Results and discussion

3.1. Powders characterization

The morphology of the as-received titanium hydride powders isdepicted in Fig. 1. The powders have an irregular shape and most ofthem have sizes of 1–40 lm. Fig. 2 displays the XRD pattern of theas-received powders. All the peaks can be indexed to those ofTiH1.924 (JCPDS PDF No. 25-0982) with a CaF2-type face-centeredcubic (fcc) crystal structure (lattice parameters a = b = c = 4.448 Å)and it can be inferred that the powders are TiH1.924. The amountof released hydrogen can reach 3.864 wt.% of the powders, whichis the maximum amount used in the foaming process of NiTiscaffolds.

Fig. 3 shows the morphology and elemental composition of theas-received powders and powders processed by TGA in helium.There is no obvious difference between the as-received and treatedpowders, as shown in Fig. 3a and b, indicating that heat treatment

Fig. 1. SEM image of as-received titanium hydride powders.

at even 600 �C for 2 h in helium does not alter the shape and size ofthe powders. On the other hand, the EDS spectra in Fig. 3a1 and b1are quite different. Only Ti can be detected in the large particles(indicated by the red frame in Fig. 3a), implying that the as-re-ceived powders are high purity titanium hydride (hydrogen cannotbe detected by EDS). After TGA a large amount of oxygen, about 67at.% oxygen, can be detected in the large particles, as shown inFig. 3b1, suggesting that titanium in the titanium hydride powdersis easily oxidized at elevated temperature despite the continuoushelium flow. Our observations differ from some reported results[27,30,31]. The phenomenon is corroborated by the correspondingXRD pattern. As shown in Fig. 4, the powders after TGA at 600 �Care composed of TiO2 with a tetragonal rutile structure andhexagonal Ti3O according to JSPDS PDF No. 21-1276 and No. 01-073-1583, respectively. To further confirm the results, the TGAwas repeated four times under nitrogen or helium from room tem-perature to over 450 �C and similar results were obtained. All theresults indicate that Ti in the titanium hydride can be readily oxi-dized during TGA in spite of the use of a protective gas.

TEM and SAED showed that the as-received titanium hydrideparticles are crystalline. The SAED pattern suggests that the largeparticles consist of clusters of several TiH1.924 crystals with differ-ent orientations, as shown in Fig. 5a. In contrast, after TGA theSAED pattern exhibits a series of polycrystalline rings, as shownin Fig. 5b. The results are in good agreement with the XRD results.

Fig. 6 shows the morphological evolution of the powders withtemperature in air. After treatment at 650 �C the size and shapeof the large particles are similar to those of the as-measured pow-ders shown in Fig. 1, whereas the smaller particles become round,as shown in Fig. 6a. The powders sintered and aggregated at tem-peratures over 700 �C are different, as shown in Fig. 6b–d. Titaniumis formed from the titanium hydride powders, oxidized, and sin-tered together at the higher temperature, as confirmed by the cor-responding XRD patterns, to be discussed later in this paper. Fig. 7displays the evolution of the Ti and O atomic content of the pow-ders with temperature during treatment in air. The Ti concentra-tion decreases quickly while that of O increases significantlywhen the temperature is raised from 300 �C to 550 �C, and the ef-fect is particularly pronounced between 500 �C and 550 �C. In thistemperature range decomposition of the titanium hydride pow-ders is very fast and titanium is readily oxidized. In the higher tem-perature range 550 �C to 650 �C the concentrations of these twoelements are relatively constant. It can be deduced that most oftitanium hydride particles have decomposed and most of the tita-nium is oxidized at around 550 �C. In the higher temperature range

Fig. 3. (a) As-received TiH2 powders; (a1) EDS result for the particle in (a) indicated by a red frame. (b) TiH2 powders after TGA in helium from room temperature to 600 �Cwith a 2 h hold at 600 �C; (b1) EDS result for the particle in (b) indicated by the red frame.

Fig. 4. XRD pattern of titanium hydride after TGA in helium from room temperatureto 600 �C with a 2 h hold at 600 �C.

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650–800 �C the O content increases slightly, revealing that a smallamount of titanium hydride in the inner cores of the large particlescontinues to decompose and be oxidized. At temperatures higherthan 800 �C the elemental content is relatively stable, implyingthat all the titanium hydride has decomposed and oxidized. Theonly difference is the valence, which it will be discussed later inthis paper.

3.2. Thermal evolution and phase transition of titanium hydride

3.2.1. DSC analysisHeat flow during the heating process in argon was monitored by

DSC. As shown in Fig. 8, there are two predominant endothermic

peaks, p1 from 310 �C to 475 �C and p4 from 540 �C to 700 �C. Sincedecomposition of titanium hydride requires external energy, thereare obvious decomposition reactions involving two different phasetransitions in these two temperature ranges, especially near 460 �Cand 580 �C, otherwise these two should merge together. The firstdecomposition commences at about 310 �C. As the temperatureis increased from 310 �C to 440 �C the reaction proceeds slowly,but accelerates in the next temperature range 440–475 �C, indicat-ing that the first stage of titanium hydride decomposition is almostcomplete at 475 �C. In the next range 475–540 �C there is a smallendothermic reaction, as indicated by peaks p2 and p3 in Fig. 8.The second decomposition stage occurring in the temperaturerange 540–700 �C requires the supply of a large amount of externalenergy. In addition to the predominantly endothermic reaction,there is a small exothermic reaction. For example, between700 �C and 750 �C an exothermic peak, p5, in Fig. 8 correspondsto a minor oxidation process, even under argon.

3.2.2. TGAThe TGA results acquired for the as-received titanium hydride

powders in Fig. 9 are quite different from the thermogravimetrycurve obtained in helium [30,31]. It can be observed that theweight increases from 250 �C to 1000 �C, and this is relatively sig-nificant between 560 �C and 950 �C. The total weight loss fromTiH1.924 can be as high as 3.864 wt.%, whereas the practical weightincrease is about 58 wt.% at 1000 �C. This is close to the oxygencontent increase of 64 wt.%, assuming that TiH1.924 is completelyconverted to TiO2. Therefore, it is believed that Ti formed fromTiH1.924 is oxidized almost simultaneously during the course ofthermogravimetry and oxidation to TiO2 is complete at 1000 �C.Since the weight loss due to hydrogen is much smaller than theweight increase because of oxygen, hydrogen release cannot be de-tected during this thermogravimetric process. The derivative ther-mogravimetry (DTG) curve is obtained from the first derivative ofthe thermogravimetry curve and the peaks correspond to the fast-est rate in the relevant range. The DTG curve in Fig. 9 reveals that

Fig. 5. TEM images and SAED patterns of titanium hydride powders: (a) as-received powders; (b) after treatment with TGA in helium from room temperature to 600 �C with a2 h hold at 600 �C.

Fig. 6. Morphologies of titanium hydride powders treated in air at elevated temperature of: (a) 650 �C; (b) 700 �C; (c) 900 �C; and (d) 1000 �C.

S. Wu et al. / Acta Biomaterialia 7 (2011) 1387–1397 1391

oxidation occurs at about 250 �C and is relatively rapid at 290 �Cand 480 �C. It possibly involves a slow and stable phase transitionprocess in this early stage. The subsequent reactions occur quicklyat about 560 �C and 600 �C and it can be deduced that decomposi-tion of titanium hydride is fastest at these two points. The XRD pat-terns and real time monitoring of the pressure in the high vacuum

CVD system confirm this phase transition process and hydrogen re-lease rate, respectively (to be discussed below). In the temperaturerange 660–1000 �C formation of titanium oxide proceeds withoutdecomposition of titanium hydride and the highest rate is at855 �C. The XRD patterns obtained from the titanium hydridecorroborate this hypothesis. The TGA results acquired from the

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1392 S. Wu et al. / Acta Biomaterialia 7 (2011) 1387–1397

titanium hydride under different atmospheres are presented inFig. 10. Similarly to air, nitrogen and helium cannot protect thepowders from oxidation during TGA. However, the thermogravi-metry curve in argon shows obvious weight losses from 480 �C to600 �C, although in the two ranges from 300 �C to 480 �C and600 �C to 1000 �C, there are weight increases of 0.5% and 3.6%,respectively, which are far lower than those under other atmo-spheres. It is believed that helium and nitrogen are unable to re-move the oxygen attached to the surface of the powders due totheir light masses and, in this regard, argon is a better gas to pro-tect titanium hydride from oxidation. Our TGA results obtained inair indicate an oxidation trend similar to that reported previously[31], but different from that reported by Malachevsky and D’Ovid-io, who observed decomposition without oxidation during heatingfrom 300 �C to 800 �C in air [27]. Malachevsky et al. ascribed thelow oxidation of titanium hydrides to the non-isothermal heatingprocess and hydrogen diffusion through the increasing superficialtitanium oxide layer in their study, but Kennedy et al. believed thatformation of an oxyhydride phase on the outer surface of titaniumhydride powders induced less oxidation when titanium hydridewas heated in air [27,29]. In comparison, our real-time monitoringin vacuum discloses a rapid hydrogen release rate which cannot beachieved by hydrogen diffusion. In addition, the XRD patterns can-not detect any traces of an oxyhydride phase. Therefore, titaniumhydrides undergo quick oxidation when exposed to air.

3.2.3. X-ray diffractometryThe phase transition during heating in air is characterized by

XRD. Fig. 11 shows the phase transition of titanium hydride asthe temperature is increased from 300 �C to 1000 �C. The fullXRD pattern in Fig. 11a shows the peak change and shift specifi-cally at 500 �C and 550 �C. At 550 �C tetragonal titanium dioxide(rutile) forms. The variation from room temperature to 600 �C inFig. 11b is indicated by a narrow scan between 33� and 42�. The(1 1 1) peak of the as-received TiH1.924 gradually shifts to a largerangle as the temperature is increased from 300 �C to 600 �C. At500 �C TiH1.5 forms (lattice parameters a = b = c = 4.402 Å, JCPDSPDF No. 01-078-2216). This reveals that TiH1.924 is gradually con-verted into TiH1.5 with the same crystalline structure as the parentphase, and that decomposition occurs rapidly in air between 500 �Cand 600 �C. This is in good agreement with the TGA results in Fig. 9.At 550 �C no TiH1.924 can be detected, implying completion of the

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S. Wu et al. / Acta Biomaterialia 7 (2011) 1387–1397 1393

transition from TiH1.924 to TiH1.5. During this process oxidation alsooccurs simultaneously, but it cannot be detected from the XRD pat-terns below 500 �C because only a trace amount of titanium oxideforms at the low temperature, but it is indicated by the mass in-crease in Fig. 9. As the temperature increases a small Ti3O (�1�12)peak is detected at 500 �C, followed by emergence of the TiO2 (ru-tile) (1 1 0) peak at 550 �C.

The phase evolution in the higher temperature range is shownby the narrow scan in Fig. 11c. From 500 �C to 650 �C the TiH1.5

peak decreases, disappearing at 650 �C, revealing decompositionof TiH1.5 in this temperature range with completion at 650 �C. Notitanium can be detected, indicating that Ti is instantly oxidized.This is in agreement with the DTG results between 500 �C and660 �C in Fig. 9. The Ti3O peaks are stable between 550 �C and700 �C, implying the formation of a large amount of Ti3O. The sig-nal decreases at 800 �C and vanishes almost completely at 900 �C.It can be concluded that Ti3O is converted into TiO2 gradually from700 �C to 900 �C. The results are consistent with the DTG results inFig. 9. As a result, the amount of TiO2 increases gradually from550 �C to 1000 �C.

Fig. 12 displays the decomposition process of titanium hydridein the high vacuum CVD system under a small argon flow. The sur-vey scan in Fig. 12a does not show the formation of titanium oxideduring heating compared with that in air (Fig. 11a) and there is noobvious change in the XRD patterns between 450 �C and 550 �C. Tofurther investigate the phase evolution, the XRD patterns acquiredat lower temperature are shown as narrow scans between 33.5�and 41.5� in Fig. 12b. In comparison with the as-received titaniumhydride at room temperature, the main peak (1 1 1) shifts to ahigher angle until 400 �C, indicating the formation of TiH1.5, whichreaches a maximum concentration near this temperature. In addi-tion to TiH1.5, a small amount of TiH1.7 (JCPDS PDF No. 01-078-2215) is formed at 400 �C. Hence, TiH1.924 decomposes predomi-nantly into TiH1.5 and a small amount of TiH1.7 starting at 350 �Cand continuing at 400 �C and higher under vacuum. TiH1.7 almostvanishes while a new phase a-Ti emerges at 450 �C. The TiH1.5 peakintensity diminishes and no signal can be detected at 600 �C inFig. 12c, while the signal from a-Ti increases significantly, reachinga maximum at 600 �C suggesting that TiH1.5 decomposes into a-Ti(JCPDS PDF No. 03-065-9622). It is believed that decompositionfrom TiH1.5 to a-Ti proceeds between 400 �C and 600 �C. Thisobservation is in line with the trend between 480 �C and 600 �C de-rived from the DSC and thermogravimetric analysis in argon, asshown in Figs. 8 and 10, respectively. Thermogravimetry in argoncannot directly detect the decomposition below 480 �C and over600 �C due to oxidation, while XRD in air could detect the distinctsignal of TiH1.5 until 500 �C, as shown in Fig. 11b. This means thatoxidation could alter the decomposition from TiH1.924 to TiH1.5, aswell as TiH1.5 to a-Ti. In addition, the conversion from TiH1.924 toTiH1.7 cannot be detected by XRD in air.

3.3. Hydrogen release evaluation

Hydrogen released from titanium hydride can be directly andprecisely detected and evaluated by real time monitoring of thepressure in the vacuum system. As shown in Fig. 13a, the pressurevaries with temperature. At 350 �C the hydrogen pressure in-creases slightly, as illustrated by the magnified graph. Accordingto the ideal gas law: DnRT ¼ DpV , where Dn and Dp are the varia-tions in gas molar weight and pressure in the system, respectively,T and V are the absolute temperature and volume of the system,respectively, and R is the gas constant of 8.314 J K�1 mol�1. In a sta-ble system, i.e. constant temperature and volume, the variation inDp must be induced by the gas molar content. Therefore, it can beinferred that hydrogen is continuously released from titanium

Fig. 12. Phase transition of titanium hydride heated under high vacuum for 30 minat different temperature measured by XRD: (a) full patterns between 20� and 80�;(b) narrow scan of XRD patterns from room temperature to 450 �C between 33.5�and 41.5�; (c) narrow scan of XRD patterns from 480 �C to 1000 �C between 33.5�and 41.5�. *, TiH1.924; , TiH1.5; h, TiH1.7; and �, a-Ti.

Fig. 13. Pressure variation induced by hydrogen release from TiH1.924 undervacuum: (a) as a function of holding time at different temperatures and (b) duringheating. The values shown by the dotted lines in (a) are beyond the range of thevacuometer.

1394 S. Wu et al. / Acta Biomaterialia 7 (2011) 1387–1397

hydride during the holding period at 350 �C. The XRD pattern inFig. 12b discloses that the (1 1 1) peak in TiH1.924 shifts towards(1 1 1) in TiH1.5. As the temperature is increased, for example at390 �C the pressure increase can be more easily detected and it isalmost constant during the last 10 min of the holding period. Obvi-ously, more hydrogen is released at higher temperatures becauseof the supply of energy for the decomposition of TiH1.924. Theapparent increment in endothermic heat flow can be clearly ob-served from 360 �C, as shown in the DSC curve in Fig. 8. InFig. 12b the (1 1 1) peak shifts closer to that of TiH1.5. Meanwhile,a trace of the (1 1 1) peak can be detected from TiH1.7 and so it isbelieved that hydrogen is released during the transition fromTiH1.924 to TiH1.5 and a smaller amount from TiH1.924 to TiH1.7.The pressure stability during the last 10 min indicates that hydro-gen release reaches an equilibrium at 390 �C. According to the XRDresults in Fig. 12b conversion of TiH1.924 into TiH1.5 and TiH1.7 is al-most complete at 400 �C and, consequently, the maximum amountof released hydrogen could reach only 0.85 wt.%. Therefore,although the pressure at 425 �C increases faster than at 390 �C inthis early stage, the equilibrium pressure increases only slightly.As the temperature is further increased to 480 �C the rate of

Fig. 14. Cross-section morphology of NiTi foams: (a) control foam without TiH1.924;(b) foam with TiH1.924 prepared by process 1; (c) foam with TiH1.924 prepared byprocess 2.

S. Wu et al. / Acta Biomaterialia 7 (2011) 1387–1397 1395

pressure increase is higher and the equilibrium pressure is alsohigher, as illustrated in the inset in Fig. 13a. This is because a largeamount of TiH1.5 is converted into a-Ti and TiH1.7 is almost com-pletely transformed into a-Ti, according to Fig. 12c. Hence, a largeamount of hydrogen is released. Except at 350 �C, the pressurereaches equilibrium early, from 390 �C to 480 �C because a hightemperature provides the needed energy for the rapid decomposi-tion of titanium hydride. Consequently, decomposition achievesstability in the early stage and the pressure decreases as theamount of hydrogen released decreases. The hydrogen releasebehavior at 550 �C and 630 �C is different from that at the lowertemperature. At 550 �C the pressure increases precipitously be-yond the maximum range of the vacuometer within 6 min. Afterabout 6 min the pressure falls back into the allowable range ofthe vacuometer and continues to decrease significantly. The resultsshow that decomposition is fast and releases a large amount ofhydrogen in a short time. The XRD pattern in Fig. 12c reveals thatonly a small amount of TiH1.5 remains. Below 630 �C decomposi-tion reaches the highest rate but the pressure subsequently de-creases to a level lower than that at 550 �C after 10 min at thistemperature, suggesting that decomposition is possibly complete.This is in agreement with the XRD data in Fig. 12c, in which onlya-Ti can be observed at 600 �C and all titanium hydride has beenconverted into a-Ti.

Fig. 13b shows the pressure change due to released hydrogenin the vacuum system heated to 1000 �C and 630 �C. In the for-mer system the pressure increases slightly in the temperaturerange 250–550 �C, but increases rapidly after about 550 �C. Thetrend is similar in the latter system, with the only differencebeing that the inflexion temperature changed to �580 �C. Thisdifference is due to the heater reducing heating close to the des-ignated temperature of 630 �C in order to reach this temperatureprecisely. In both cases the pressure very quickly exceeds therange of the vacuometer when the temperature is >600 �C, indi-cating maximum hydrogen release at this temperature. This is inagreement with the maximum endothermic heat flow at 590 �Cshown in Fig. 8. For the first system (heated to 1000 �C) the pres-sure is in the instrumental range and decreases significantly at900 �C.

3.4. Optimizing foaming of the NiTi scaffold

Based on our understanding of hydrogen release from and thephase transitions of titanium hydride in air and argon and undervacuum we are able to optimize the foaming process describedabove. The effects of hydrogen release on foaming are studiedusing the cross-sectional morphology in Fig. 14. The control foamwithout TiH1.924 exhibits a smaller porosity of about 18 vol.%caused by argon during the early stage of the CF-HIP process.Most of the pores are almost round and distributed non-uniformly, as shown in Fig. 14a. According to Fig. 14b and chydrogen released from the titanium hydride significantly en-hances the porosity of NiTi foams. In comparison with the controlfoam, the NiTi foam produced from a mixture of Ni and TiH1.924

possesses a much larger porosity of over 60 vol.%, as well as largepores. The measured open porosities of the two types of foamsexceed 80% according to the standard protocol [37], indicatingthat the released hydrogen significantly enhances pore inter-connection.

The foams with TiH1.924 produced by the two different pro-cesses have different porous structures. In Fig. 14b the foam pro-duced by process 1 shows a uniform pore distribution and mostpores are nearly round with a relatively homogeneous size of100–250 lm. Only a few merged pores have sizes of 250–500 lmand the structure favors cell and tissue in-growth, as well as mass(i.e. liquid and blood) transport [38,39]. The pores in the structure

produced by process 2 have an irregular shape and distribution.The size of most pores is about 30–150 lm and only a few poresare larger than 300 lm. The former size is much better than thelatter with regard to cell and tissue in-growth and load bearingin tissue repair or construction. The difference is caused by the dif-ferent hydrogen release behaviors in the two processes. As men-tioned, hydrogen release is associated with the phase transitionsof titanium hydride. Although this occurs throughout the heatingprocess from about 300 �C to 700 �C according to the thermal

1396 S. Wu et al. / Acta Biomaterialia 7 (2011) 1387–1397

analysis and XRD analysis, the main phase transition occurs via thefollowing two stages: (1) TiH1.924 ? TiH1.5/TiH1.7 from 300 �C toabout 400 �C and (2) TiH1.5/TiH1.7 ? a-Ti from 400 �C to about600 �C. Since minor oxidation is unavoidable during the CF-HIPprocess [22], hydrogen release can be deferred to a higher temper-ature under high pressure argon during CF-HIP. In the first stagethe rate of hydrogen release is slow, but increases precipitouslyin the second stage. According to Wilkinson and Ashby’s expansionmodel of single closed pores [40] the pores expand when the inter-nal pressure rises sharply over a short time period, inducing a mis-match between the inside and outside pressures. The pores canbecome irregular or crack at the weakest point if the strength ofthe pore walls is not uniform. In process 1 hydrogen is releasedin a stepwise fashion at 425 �C, 480 �C, 500 �C, 550 �C, and600 �C. According to Fig. 13a, early holding reduces the hydrogenrelease rate during subsequent holding. As a result, process 1 in-duces a gradual change in pressure and so the pores are nearlyround and, as they become bigger, they touch adjacent ones, asillustrated in Fig. 14b. In comparison, direct heating in process 2leads to hydrogen release as depicted in Fig. 13b, showing that attemperatures below 550 �C hydrogen release is very slow, but sta-ble. The pores do not expand in this stage. However, when the tem-perature exceeds 500 �C rapid release of hydrogen inducesdeformation and deflation of the pores, giving rise to pores of smallsize and irregular shape, as shown in Fig. 14c.

Fig. 15. SEM images of cell morphologies on NiTi scaffolds foamed by hydrogenreleased from titanium hydride after culture for 8 days: (a) on the entire surface;and (b) on the surface of the exposed pore.

3.5. Cytocompatibility of NiTi scaffolds

Porous metal foams such as titanium, tantalum, NiTi, and tita-nium alloys are attracting more attention from biomaterials scien-tists and orthopedists due to their porous structure, which allowstissue in-growth and long-term fixation of implants, in additionto their good mechanical properties and lower Young’s modulus[38,41–46]. Although many factors influence the behavior of cellsand the in-growth of tissues into porous meal implants, for exam-ple the biomechanical properties [47] and motion of the implant,as well as the presence of gaps between the bone and implant[42], the pore size in the implant plays a critical role in cell migra-tion and tissue in-growth. It is recognized that pore sizes of over50 lm are necessary in order for osteoblasts to traverse into theporous channels [22,38,48]. Consequently, the foam produced byprocess 1 with a pore size of 100–250 lm was selected for furthercytotoxicity evaluation. As shown in Fig. 15a, bone cells can attachand proliferate on the entire surface of the NiTi foams. Some smal-ler pores are covered by the bridged osteoblasts. It can be observedfrom Fig. 15b that osteoblasts can easily grow inside the exposedpores, including the side surfaces and bottom. The preliminary cellculture reveals that the NiTi foams fabricated by process 1 havegood cytocompatibility and the corresponding porous structure fa-vors the in-growth of osteoblasts.

4. Conclusion

We have systematically investigated hydrogen release fromtitanium hydride under air, argon, and vacuum using thermalanalysis and real time monitoring of the pressure under vacuum.Hydrogen release occurs in two stages: (1) TiH1.924 ? TiH1.5/TiH1.7 between 300 �C and 400 �C and (2) TiH1.5/TiH1.7 ? a-TiTiH1.5/TiH1.7 ? a-Ti between 400 �C and 600 �C. During the heatingprocess the hydrogen release rate is slow initially while the decom-position rate of titanium hydride increases quickly above 550 �C.The amount of released hydrogen is influenced by the temperatureand holding time. A stepwise heating process at 425 �C, 480 �C,500 �C, 550 �C, and 600 �C during CF-HIP significantly improvesthe porous structure in the NiTi scaffolds. A preliminary cell culturetest reveals that the NiTi foams produced by this process do not in-duce obvious toxicity and the pore size is suitable for the in-growth of osteoblasts. In addition to orthopedic applications ofNiTi scaffolds, hydrogen storage and release are of increasing inter-est from the perspective of energy research and environmental sci-ence [49–52]. The fundamental understanding and technologydescribed here can be extended to other hydrogen storage materi-als, the foaming of industrial metals, and other types of biomedicalscaffolds.

Acknowledgements

This work was jointly supported by City University of HongKong Strategic Research Grant (SRG) No. 7008009, National NaturalScience Foundation of China Grant No. 50901032, Ministry of Edu-cation Specialized Research Foundation for University DoctoralProgram Grant No. 20094208120003, and Hubei Provincial NaturalScience Foundation Grant No. 2009CBD359. S.L. Wu thanks Dr KailiZhang (City University of Hong Kong) for assistance with the TGanalysis.

Appendix. Figures with essential colour discrimination

Certain figures in this article, particularly Figs. 3, 4, 7–13, aredifficult to interpret in black and white. The full colour images

S. Wu et al. / Acta Biomaterialia 7 (2011) 1387–1397 1397

can be found in the on-line version, at doi:10.1016/j.actbio.2010.10.008.

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